The substrate material for this experiment is quenched and tempered AISI 1045 steel with dimensions of 50 mm × 50 mm × 15 mm. Prior to the experiment, the substrate material underwent grinding and matte treatment to reduce reflection and maximise energy absorption. In previous studies, the authors determined the optimal mass fractions of Fe‒Ni‒Ti cladding powders, as shown in Table 1.

The prepared deposition powder undergoes ethanol-protected ball milling in an XQM-2 L vertical planetary ball mill using agate grinding balls with diameters of 6, 10, and 20 mm. Among them, large balls constitute 20% of the total weight, medium balls 40%, and small balls 40%. The ball milling duration is 12 h, the grinding ball-to-deposition material ratio is 3:1, and the milling speed is 250 r/min. The particle size distribution of the milled cladding powder is tested using a laser particle size analyser; the results are shown in Figure 1. The surface density of the melted powder after ball milling measured by the porosity tester is 3.32 g/cm³, and the internal porosity is 15%. The flowability of the cladding powder is tested using a powder feeder (BTSF-2). The feeding accuracy is 99.3%, indicating the excellent flowability of the cladding powder.

The experimental setup utilises an LDM2500-60 semiconductor laser for laser deposition, accompanied by a BTSF-2 coaxial powder feeder, with Ar gas protection during the deposition process. The deposition path is controlled by the numerical control system of the laser processing machine. Single-track, single-layer laser deposition experiments are conducted on the surface of grade AISI 1045 steel using coaxial powder feeding. The experimental process parameters are listed in Table 2. The working principle of laser deposition is shown in Figure 2.

The laser energy density can be more clearly expressed as the thermal input per unit area, defined by Eq. 1 as follows: ρ = P A = P V B (1) where ρ represents the laser energy density (W/mm²); P represents the laser power (W); A denotes the clad area (mm²); V represents the scanning speed (mm/min); and B represents the scanning width (mm). The calculation according to Eq. 1 yields the laser power and corresponding energy density, as presented in Table 3.

The laser-deposited samples are cut into specimens measuring 5 mm × 4 mm × 4 mm using a metallographic cutting machine (DK733-A). The surface of the deposited coating is cleaned, and the cross section is subjected to friction and polished. A 20% nitric acid alcohol solution is used for etching, with an etching time of 10 s. The surface pores and cracks of the deposited coating are observed under an optical microscope (Vert. A1), and the porosity is quantified using ImageJ software. The pore area of the metallographic structure is measured using ImageJ software, and the porosity of the cladding layer is calculated using the following formula: Porosity = Pore Area/Total Area.

The density is measured using the Archimedean displacement method. The density of the composite sample is measured according to Archimedes’ principle. The calculation of actual density can be found in Eq. 2. ρ = m 1 m 2 × ρ 1 (2) where m₁ is the mass of the deposition layer in air (g); m₂ is the mass of the deposition layer in liquid (g); and ρ₁ is the density of distilled water (1 g/cm³). Five sets of samples are measured, and the average value is taken, ensuring that the samples are dry and that the surfaces are clean during the measurement process. The density η of the sample is the ratio of the actual density ρ to the theoretical density ρ,’ defined by Eq. 3 as follows: η = ρ ρ ′ × 100 % (3)

Additionally, the deposited samples were sectioned for observation of the coating microstructure using a scanning electron microscope (Tesan VEGAII lMH). The deposited samples are prepared as sample blocks with dimensions of 10 mm × 15 mm × 2 mm, and the phase composition and residual stresses of the cladding layer are analysed using XRD (Bruker D8 ADVANCE). The hardness of the deposition layer is measured using an HV-1000IS microhardness tester. The friction coefficient of the cladding layer is measured using a multifunctional material surface performance tester (MFT-4000) under a load of 50 N for 20 min with a 3 mm diameter abrasive ball and a rotation speed of 150 r/min. Use a universal testing machine (SHT-4605) to test the tensile strength and elongation of the samples, with a tensile speed of 1 mm/min. Each group of samples is tested 5 times, and the arithmetic mean of the tensile test results is taken. The cladding layer is subjected to sliding wear testing using an M-2000 friction and wear testing machine under a load of 250 N, a spindle speed of 200 r/min, and a friction time of 30 min. The samples are weighed before and after wear using an electronic analytical balance (J323A) with an accuracy of 1 mg, and the percentage of mass loss relative to the original mass is used to characterise the wear performance. The worn morphology of the cladding layer is captured using a scanning electron microscope.

XRD is primarily used to measure residual stress near the surface of materials. For most applications, understanding the stress state of the material’s surface is crucial. Although the penetration depth of XRD is limited, it is usually sufficient to cover the stress regions of interest in many engineering applications. Figure 3 shows the significant influence of the laser energy density on the residual stress of the deposition layer. As the energy density increases, the residual stress in the upper part of the deposition layer gradually decreases. However, as it approaches the substrate, the residual stress undergoes rapid growth, reaching its maximum at the metallurgical bond between the coating and the substrate. The maximum residual stress in the upper region of the deposition layer is 220 MPa, corresponding to a laser energy density of 83.3 J/mm². At the interface between the deposition layer and the substrate, the maximum residual stress reaches 270 MPa, corresponding to a laser energy density of 111 J/mm². When the laser energy density is 97.2 J/mm², the residual stress in the deposition layer is minimised.

During laser deposition, the cladding layer and substrate material undergo plastic deformation under the influence of the laser. The laser energy density, as a key factor determining the morphology of the cladding layer and the trend of crack formation, is directly linked to the formation and distribution of residual stress in the cladding layer (Luu et al., 2022). An increase in the laser energy density promotes the uniformity and density of the cladding layer, thereby reducing the thermal stresses in the deposition region. At the same time, an increase in the laser energy density also results in a larger temperature gradient between the cladding layer and the substrate material, consequently leading to greater thermal and structural stresses (Hector and Hetnarski, 1996). These stresses cannot be completely relieved during the cooling process of the cladding layer; thus, these stresses remain in the interior of the cladding layer, forming residual stress.

Figure 4 illustrates the microstructure and structural defects of the deposition coatings under different laser energy densities. The data show that the microstructure of the composite coating contains dendrites and equiaxed grains. Figure 4F shows the trend of the pore percentage for the five sets of samples. The pore percentage of the coating decreases initially and then rapidly increases with increasing energy density. When the laser energy density ranges from 83.3 J/mm² to 97.2 J/mm², the percentage of pores in the metallographic structure of the deposition layer gradually decreases from 2.9% to 1.2%. At a laser energy density of 97.2 J/mm², the percentage of pores in the Fe‒Ni‒Ti coating reaches its minimum value of 1.2%, and few pores are observed in the metallographic structure. As the energy density continues to increase, large pores appear in the metallographic structure of the deposition layer, and the pore percentage increases from 1.2% to 3.7%. A relatively low laser energy density results in limited heating of the deposition powder and uneven heating, leading to partially unmelted debris in the molten pool, causing the formation of pore defects in the microstructure (Lu et al., 2022).

A lower laser energy density results in less heating of the deposited powder and uneven heating, leading to the presence of partially unmelted particles in the melt pool, causing higher residual stresses in the deposited coating and resulting in the formation of porosity defects in the structure. Conversely, when the laser energy density is too high, the fluidity of the metal melt significantly increases, resulting in the splattering of metal droplets in the melt pool, exacerbating the instability of the deposition region, and leading to the formation of large pores in the structure (Manoj et al., 2023). From the perspective of porosity, the optimal energy density for the Fe‒Ni‒Ti deposition layer is 97.2 J/mm².

When the deposition powder is irradiated with a laser beam, there is an extremely high temperature slope in the molten pool, leading to a significant surface tension gradient. This results in turbulent thermocapillary flow in the molten pool, known as the Marangoni flow (Jiang et al., 2020). The stronger the laser absorption rate is, the greater the temperature of the mixed melt in the molten pool, causing an increase in Marangoni flow (Alizadeh-Sh et al., 2020).

Figure 5 illustrates the effect of the laser energy on the density of the deposition layer. When the scanning speed and powder feeding rate are kept constant, the density of the deposition layer increases with increasing laser energy density. When the density reaches a peak of 99.1%, it no longer increases but instead decreases to 96.5% with a further increase in the laser energy density. This phenomenon is primarily based on the variation in the flow state of molten metal in the molten pool due to changes in energy density during the laser deposition process. Marangoni convection is adopted as an indicator to measure the strength of liquid metal convection in the molten pool (DebRoy et al., 2018) as shown in Eq. 4. M a = d γ d T − L Δ T μ a (4) where Ma represents the degree of Marangoni convection, μ is the viscosity of the molten metal, α is the heating diffusion coefficient of the alloy, L is the width of the deposited coating, ΔT is the distinction between the highest temperature and the solidus temperature of the alloy, and d γ d T represents the temperature sensitivity to the exterior pressure.

According to the laws of physics, there is a negative correlation between the kinematic viscosity of the fused metal or liquid metal in the laser molten pool and the ΔT value as shown in Eq. 5. Δ T = 2 A E k k t h τ p π (5) where A represents the laser absorption rate, E is the laser energy density, k is the heat conductivity, kth is the heat diffusivity, and τ_p is the laser radiation time (Zhang et al., 2022).

From the above formulas, it can be concluded that a lower energy density can result in the presence of partially unmelted particles in the molten pool, leading to a lower alloy structural density. Appropriately increasing the laser energy density can achieve a suitable ΔT, allowing for sufficient melting of the alloy powder during the forming process and improving the density. However, an excessively high laser energy density can result in an excessively high maximum temperature in the molten pool, leading to a high ΔT. An excessive ΔT intensifies the fluctuation of Marangoni convection, causing instability in the forming process, leading to defects such as pores and cracks, and reducing the density of the alloy.

Figure 6 shows the microstructure of the Fe‒Ni‒Ti deposition layer under an SEM at a laser energy density of 97.2 J/mm². The deposition layer mainly consists of fine dendrites with a few equiaxed grains. The grains grow towards the centre of the molten pool, perpendicular to the solid‒liquid interface, exhibiting characteristics of directional solidification and epitaxial growth. The central region of the deposition layer consists of equiaxed grains with an average size of 9 μm and clear grain boundaries. Figure 6C shows the microstructure of the deposition layer at 4,000x magnification, where in situ TiC reinforcement phases are observed uniformly distributed at the grain boundaries.

In the laser deposition layer, columnar grains mainly appear in the boundary area between the cladding layer and the substrate, i.e., the transition zone and the heat-affected zone. This is due to the significant temperature gradient in this area, which causes the grains to grow directionally. Equiaxed grains primarily appear in the central region of the cladding layer, which is the last cooling area in the molten pool. The grains have a longer growth time and are influenced mainly by convection within the melt.

Figure 7 presents the X-ray diffraction analysis results for the phases of the Fe‒Ni‒Ti deposited layers under different laser energy densities. The data show that the variation in laser energy density does not alter the phases of the overlay. Austenite is detected in the Fe‒Ni‒Ti deposited layer, along with the presence of a face-centred cubic (FCC) phase. The structure includes α-Fe, Fe₃Ni₂, and an in situ-formed TiC reinforcement phase. The reaction for the formation of the TiC phase is represented by Eq. 6 As follows: T i + C = T i C 3100 ° (6)

At an energy density of 97.2 J/mm², the nickel diffraction peak in the Fe‒Ni‒Ti deposited layer becomes lower and broader, indicating a refinement of the metal grains and a denser coating structure, which is conducive to enhancing the material’s mechanical properties, consistent with the SEM analysis results (Penelle and Baudin, 2010).

Figure 8 presents a column chart of the mechanical properties of the Fe‒Ni‒Ti deposited layers at various laser energy densities. The chart shows that with increasing energy density, both the tensile strength and extension of the deposited layer initially increase and then decrease (Yu et al., 2021). At an energy density of 83.3 J/mm², the tensile strength and extension of the deposited layer are minimal, measuring 545 MPa and 3.44%, respectively. With a density of 97.2 J/mm², the tensile strength of the deposited layer reaches its maximum value of 628 MPa. With a density of 104.2 J/mm², the elongation of the deposited layer reaches its maximum value of 5.8%. With a density of 97.2 J/mm², the microhardness of the deposited layer reaches its highest value of 518 HV, which is approximately 2.24 times greater than the hardness of the base AISI 1045 steel.

With increasing laser power density, the local temperature of the substrate gradually increases, causing the grains in the deposited layer to transform into fine equiaxed grains. The refinement of the grain size leads to an increase in the overall microstructural density. With a density of 97.2 J/mm², the TiC precipitate particles generated in situ are uniformly dispersed in the austenitic matrix, providing effective precipitation strengthening. During the tensile deformation process, TiC precipitate particles strongly impede the movement of dislocations, thereby enhancing the tensile strength of the coating. With a density of 111 J/mm², the microstructural density of the deposited layer decreases, and larger pore defects appear, resulting in a reduction in strength.

The distance from the substrate to the deposition layer is measured over a distance of 0.6–1.0 mm. The microhardness of the deposition layer in this region is tested, as shown in Figure 2. With increasing laser energy density, the microhardness initially increases and then decreases. An analysis of the cause is the following. The primary reason for this phenomenon is correlated with the uniform distribution of ceramic-reinforced TiC in the deposited layer. At a low laser energy density, the melt pool has insufficient energy and poor flowability, resulting in an uneven TiC distribution and lower microhardness. At excessively high laser energy density, the TiC particles undergo melting. This reduces the tendency for dislocation slip, leading to a decrease in the microhardness of the deposited layer. On the other side of the shield, an increase in the energy density increases the internal temperature of the deposited layer. This results in a slower cooling rate, reduced temperature gradient, decreased undercooling, and substantial growth of the internal structure. Consequently, grain coarsening occurs, leading to a decrease in the quality of the cladding layer and a decrease in hardness (Ma et al., 2014).

The variation in the friction coefficient over time for the Fe‒Ni‒Ti coatings deposited under different laser energy densities and subjected to a normal load of 50 N is illustrated in Figure 9A. The friction coefficient represents the ability of a material to reduce friction under complex working conditions. A lower friction coefficient indicates better anti-friction performance, resulting in reduced material wear. In the initial stage of frictional wear, labelled Stage C or the running-in phase, the friction coefficient gradually increases. After 5 min, it enters the stable wear stage labelled Stage D. With increasing relative friction time, debris continuously accumulates and reaches a stable state, indicating the onset of stable wear in the frictional process. Figure 9A shows that at a laser power density of 97.7 J/mm², the friction coefficient of the deposited layer exhibits the smallest fluctuation during the stable wear stage, indicating the most stable friction behaviour. Figure 9B illustrates the average friction coefficient of the deposited coating and substrate entering the stable wear stage after 10 min. The average friction coefficient of the deposited coating is consistently lower than that of AISI 1045 steel. With increasing energy density, the average friction coefficient of the Fe‒Ni‒Ti-deposited coating initially decreases and then increases (Feng et al., 2022; Wang Z. et al., 2023). At a laser power density of 97.7 J/mm², the deposited layer exhibits the lowest average friction coefficient of 0.391, which is mainly attributed to the greater hardness and compactness at this energy density.

This study investigates the wear performance of the deposited coating through the variation in wear mass loss and the percentage change in the original coating mass. The variation in the wear mass loss ratio of the Fe‒Ni‒Ti coatings deposited under different laser energy densities and subjected to a normal load of 250 N is shown in Figure 10. The wear loss of the deposited layer is consistently lower than that of the AISI 1045 steel substrate. Moreover, at a laser power density of 97.2 J/mm², the wear mass loss ratio of the deposited layer is minimal, measuring 0.08%. During the wear stage of the deposited layer, the melting of TiC particles into the layer plays a role in solid solution strengthening and precipitation strengthening, leading to an increase in the hardness of the deposited layer and a decrease in wear. The changes in the microhardness of the cladding layer reveal that the wear mass loss ratio is nearly directly proportional to the hardness distribution trend (Guo et al., 2010; Huang et al., 2022).

Figure 11 shows SEM images of the wear morphology of the Fe‒Ni‒Ti-deposited coating. In Figure 11A, corresponding to a laser power density of 97.2 J/mm², the wear morphology of the deposited layer shows mainly shallow furrows on a relatively flat surface, indicating mild abrasive wear. In Figure 11B, corresponding to a laser energy density of 111 J/mm², as the energy density increases, the furrows deepen, and the wear surface becomes uneven. Coating delamination occurs due to plastic deformation, indicating that the predominant wear mechanisms in the deposited layer are abrasive wear and fatigue wear (Zhang and Chen, 2006; Aghababaei et al., 2016). The uniformly distributed TiC particles in the coating, acting as a hard phase, impede the expansion of furrows during wear, enhancing the wear resistance of the material (Mishina and Hase, 2013; Molinari et al., 2018).