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Article

Friction Performance Analysis of WC-Reinforced IN718 Composite Material Based on SLM Process

by
Xuejin Zhao
,
Xiaoyu Zhao
,
Youfan Xu
and
Yongjun Shi
*
College of Mechanical and Electronic Engineering, China University of Petroleum, Qingdao 266580, China
*
Author to whom correspondence should be addressed.
Metals 2024, 14(12), 1361; https://doi.org/10.3390/met14121361 (registering DOI)
Submission received: 21 October 2024 / Revised: 21 November 2024 / Accepted: 22 November 2024 / Published: 29 November 2024
(This article belongs to the Special Issue Recent Developments in Laser Additive Manufacturing of Metallic Materials)
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Abstract

:
To enhance the wear resistance of nickel-based high-temperature alloys, IN718/WC composites were prepared via selective laser melting (SLM). The optimal molding process parameters of IN718/WC composites were determined using a combination of experimental and simulation methods. Based on the SEM images of the composites, a gradient transition layer was found to form between the WC particles and the IN718 matrix, and the γ′ and γ″ reinforcing phases dispersed into the matrix, providing lubrication and reducing wear during friction. The influence of WC content on the wear resistance of the composites was investigated. When the WC content was 5%, the molded parts showed optimal wear resistance, the friction coefficient fluctuated steadily, the degree of wear was low, and the amount of wear was reduced to 0.02973 mm3. The average friction coefficient and wear volume of the molded parts with similar contents decreased by 26.95% and 4.27%, respectively, compared with the pure IN718-molded parts. This study provides a case study and guidance for further optimization of the molding process for nickel matrix high-temperature composites prepared using the SLM method.

1. Introduction

Nickel-based high-temperature alloys are alloys with a nickel quality fraction > 50%; high strength; and good resistance to oxidation and gas corrosion at high temperatures, such as in the range of 650 to 1000 °C. They hold a special and important place in the field of high-temperature alloys, as 40% of modern gas turbine engines contain nickel-based high-temperature alloys [1]. In particular, the IN718 alloy is one of the most widely used in nickel-based high-temperature alloy systems, such as in gas turbine disks, rocket engines, and aerospace vehicles, due to its excellent corrosion, fatigue, and abrasion resistances at high temperatures, and good weldability [2,3].
The IN718 alloy has also been used in hardening processes. However, since the high-precision production of key parts with complex morphologies is difficult using traditional methods of rapid prototyping, selective laser melting (SLM) technology molding has emerged as an alternative method, boasting a high efficiency, manufacturing flexibility, and other advantages. The use of SLM to produce nickel-based high-temperature alloys, which involves a melting process first and then solidification via a series of phase transition processes, has a big impact on the resultant macroscopic mechanical properties. In general, nickel-based high-temperature alloys are multi-component alloys containing Al, Ti, Fe, Cr, Nb, Mo, and other elements, where the phases are mainly the γ matrix phase and the γ′, γ″, and δ enhanced phases, as well as the MC (metal carbide) [4,5,6,7,8].
The components made from IN718 often operate in harsh environments, leading to component failure during prolonged frictional wear. However, the wear resistance of the material can be improved using surface treatment techniques, such as heat treatment, coating, and surface modification, and can also be enhanced by adding ceramic particles [9,10,11]. The selection of the ceramic-reinforcing phase should be based on the application and cost of the material: structural materials are usually selected when high modulus strength and low density are required, while materials selected for high-temperature applications usually require a small coefficient of thermal expansion and good thermal conductivity. Considering the application of nickel-based high-temperature alloys in aerospace, including important parts of industrial gas turbine engines that work under high-heat environments, the choice of reinforcing phase must meet both the mechanical and thermal stability requirements. At present, the reinforcing phases added to nickel-based alloys are mainly carbide ceramics (TiC, WC, and SiC) and graphene, which are thermodynamically more stable and have good thermal expansion similar to that of nickel-based alloys, ensuring that nickel-based composites have excellent comprehensive mechanical properties [12,13,14,15,16,17].
Studies on SLM molding composites have shown that among the many ceramic-reinforcing phases, tungsten carbide (WC) has received significant attention in research and practical applications over the past decades. WC can form carbides and solid solutions with the metal matrix during SLM molding [18]. Specifically, WC-TiC-Co cemented carbide cutting tools have been widely used. WC can also be used as a modified additive in NbC-C and TaC-C ternary system carbides to reduce the sintering temperature and maintain excellent properties and can be used as an aerospace material. Due to its excellent high-temperature creep resistance, tungsten carbide (WC) has a very high potential for reinforcing SLM nickel-based alloys to improve mechanical properties [19]. Q.B. Nguyen [20] et al. showed that the WC ceramic-reinforcing phase during the SLM process can refine matrix grains, while the carbides distributed along the grain boundaries inhibit localized deformations during indentation, which ultimately improves the wear resistance of the alloy. Therefore, when used in the nickel-based high-temperature alloy-reinforcing phase, WC clearly has significant advantages and potential. In addition, several related studies have shown that the mechanical properties of ceramic-reinforced composites are correlated with the particle size, morphology, distribution, and interfacial bonding properties of the particles [21,22,23,24,25]. Wu et al. [26] established a correlation model between SLM printing process parameters and WC-reinforced iron-based composites through the surface response method. During the printing process, the W and C atoms diffuse from WC particles into the matrix, which significantly increases the strength of the matrix and reduces the plastic deformation ability of the matrix.
The SLM molding process is an instantaneous non-stationary physical process; SLM has a high degree of subcooling. Though the liquid phase of the melt pool time is short, the cost of observing the melt pool through physical experiments is significant, so a new means of observing the thermal behavior of the melt pool must be found in order to study the numerical simulations of finite elements. In order to study the three-dimensional temperature field of the laser molding melt pool, Xi et al. [27] used ANSYS 16.0 software to establish a numerical model and analyze the variation in the temperature field of the laser melt pool over time. Jia et al. [28] established a parametric temperature and stress finite element model of the metal laser solid forming (MLSF) process based on the secondary development language APDL of ANSYS, which comprehensively takes into consideration the characteristics of MLSF molding to increase the accuracy of the simulation and reduce the number of mesh cells. The laser power in the molding process mainly affects the energy density absorbed by the powder and the thermal history of different layers, which results in differences in the organization, mechanical properties, and accumulation of thermal stresses in different ways.
In this study, we mainly investigate the molding process of WC-enhanced IN718 composites, and the effects of different laser powers and scanning speeds on the molding quality via simulation and verification with experiments. The effects of the addition of WC ceramic-reinforcing particles to the composites on the molding quality, microstructure, and phase of the molded parts are systematically analyzed, and the mechanism behind how the added WC strengthens the IN718/WC composites is investigated.

2. Materials and Methods

2.1. Material Pre-Processing

The matrix material used in this study is the spherical 300 mesh IN718 alloy powder produced using the aerosolization method of Jiangsu Willary New Material Technology Co., Ltd. (Xuzhou, China) with particle sizes of 15 to 53 microns; the main chemical composition of the IN718 alloy is shown in Table 1. The basic method of producing an atomization powder is to use high-speed airflow to impact the liquid metal flow; to convert the kinetic energy of the gas into the surface energy of the melted metal through collision; to break the molten metal flow into small droplets; and then, to rapidly cool and solidify the liquid in an airflow atmosphere to form a powder. The atomized powder used in this study has the characteristics of a small particle size, high sphericity, and low oxygen content.
From the XRD spectra of the pure IN718 alloy in Figure 1, we can see obvious {111}, {200}, and {220} crystal diffraction peaks, which were determined to be the Ni-γ phase; high purity of the powder; and no other impurity phases. In addition, the γ′ phase with the chemical formula Ni3(Al, Ti, Nb) and the γ″ phase with the chemical formula Ni3Nb formed by the Ni, Al, Ti, and other elements were determined from the spectra.
To enhance the mechanical properties of the IN718 nickel-based high-temperature alloy, irregularly shaped WC ceramic particles with particle sizes of 5 to 15 µm produced by the same company were used in the reinforcing phase; its microstructure is shown in Figure 2b. In order to ensure sufficient mixing of the two powders, IN718/WC powder with a 1% mass fraction of the WC-reinforcing phase was mixed in a vertical planetary ball mill from Changsha DEK Instruments and Equipment, and the ball milling process was run at 200 rpm, with a 15 min shutdown per hour of ball milling for natural cooling, for a total of 4 h of ball milling. Due to the low proportion of WC in the mixed powder, the Hall flow rate of the mixed powder was based on the flowability of pure IN718, which is 25 ± 1 S/50 g. Argon gas was used as a protective gas during the ball milling process to prevent unwanted oxidation. After ball milling, the WC ceramic-reinforcing phase in the composite powder was uniformly distributed around the IN718 powder; some of the WC ceramic particles were broken during the ball milling process, and the broken WC ceramic particles were attached to the surface of the IN718 powder. The ball-milled IN718/WC composite powder was put into a vacuum-drying oven for 4 h at 60 °C. Following this drying process, the dried powder mixture did not show any agglomeration and the powder fluidity improved, which made it more uniform when spread with a scraper; uneven spreading did not occur, likely due to the local agglomeration of the powder.

2.2. SLM Equipment

The IN718/WC composite material molding experiments were carried out using the FS271M equipment of Huashu Hi-Tech (Changsha, China), as shown in Figure 3; a 500 W fiber laser; and a high-precision scanning galvanometer, with the wavelength of the laser beam being 1064 nm and the adjustable laser diameter ranging from 0.08 mm to 0.2 mm. Nitrogen was selected as the protective gas in the experimental process. The substrate was preheated at 100 degrees Celsius to reduce the initial temperature gradient, and sintering was started when the oxygen content of the molding environment was reduced to less than 0.2% to prevent oxidative deterioration of the powder during molding. The scanning strategy has a significant impact on the performance of the molded parts. To avoid the influence of anisotropy of the molded parts on the experimental results, the scanning strategy in this study was set to serpentine scanning with an interlayer rotation angle of 67°; the scanning schematic is shown below in Figure 4.

2.3. Microstructure Characterization and Mechanical Property Testing

The formed IN718/WC composites were cut off from the SLM substrate via EDM using a wire cutter, and the molded parts were wet-sanded with 320-, 800-, 1500-, and 3000-grit abrasive paper at room temperature in sequence, then polished to a mirror surface using a metal metallographic grinding and polishing machine, cleaned with anhydrous ethanol, and dried in hot air.
Metallographic characterization also required the use of a homemade aqua regia etching solution (HNO3:HCI = 1:3) on the surface of the molded parts to wipe away any corrosion, wiping for 30 s or so; quick cleaning of the corrosion surface with anhydrous ethanol; and then drying of the microstructure for observation using a metallographic microscope. The physical examination program was used to examine the physical composition of the IN718/WC molded parts using the XRD technique to determine the physical composition.
The densification experiments were measured using the Archimedes drainage method. The density of the theoretical molding material was defined according to the material properties introduced in the previous section; the mass of the molded part in air was weighed as m0; the molded part was lifted up via thin wire winding, so that the molded part was slowly immersed into the liquid container without direct contact with the weighing container; the difference in weight before and after immersing the molded part into the liquid-containing container was weighed as ∆ m ; and the densities of the SLM-molded parts were calculated using the following formula:
ρ = m 0 ρ L ∆ m ρ s
The density of the molded parts is expressed as ρ in units of 1. The density of a liquid medium is expressed as ρ L in k g / m 3 . ρ s denotes the theoretical density of the material in k g / m 3 . m 0 denotes the actual mass of the molded part in k g . The difference in mass of the container before and after the test is ∆ m in k g .
A friction test was carried out using the CFT-1 friction and wear testing machine to conduct linear reciprocating friction and wear tests, and the friction partner was a 4 mm diameter silicon nitride ceramic ball. The test parameters of the friction experiment were room temperature, the loading mode as the weight loading, a loading mass of 500 g, a loading time of 30 min, a reciprocating frequency of the friction experiment of 5 Hz, and a friction slip distance of 5 mm. The friction coefficient was recorded throughout the friction experiment. After the friction experiment, the wear morphology was observed, the wear area morphology was plotted, and the experimental wear amount was calculated according to the morphology.

3. Results

3.1. Research on Molding Processes

The mechanism whereby the scanning speed and laser power influence the SLM-molded parts is actually the same mechanism whereby the laser energy absorbed per unit volume of powder influences the molding process, and a change in either the scanning speed or laser power is essentially a change in the laser energy density, which thus influences the molding quality. Therefore, the influence of laser energy density on the molding of IN718/WC composites will be discussed from the perspective of laser energy density. Laser energy density is defined as follows:
e = P v t d
Since the thickness of the laying powder and the spot diameter are kept constant in this experiment, t × d is a constant value, so the study of body energy density can be simplified to the study of laser line energy density:
e L = P v
where e denotes the laser body energy density in units of J / m 3 , P denotes the laser power in W , v is the scanning speed in m / s ,   t denotes the thickness of powder spread in m , and d denotes the spot diameter in m . Figure 5 shows the variation in density and microhardness of the SLM-molded parts with the laser line energy density.
When the energy density is low, the melt pool temperature is overly low; the absorbed heat is insufficient to melt the IN718/WC composite powder completely; the melt pool morphology is poor; the microhardness and molding densities are low; the width of the melt pool exceeds the scanning interval, as indicated by the green color in the simulation results; and the scanning interval is not reached, as indicated by the lack of color in the simulation results. From Figure 5, we can see that when the energy density is overly low, the depth of the melt pool is less than the thickness of the laying powder, the maximum width is less than the scanning interval, the overlap and metallurgical bonding between the layers of the molded parts is poor, large amounts of unmolten powder are present, the molded parts show a large number of irregular pores, the powder is sticky, and the metallurgical phase of the molded parts is porous. As shown in Figure 6 and Figure 7, the low energy density caused by a low laser power or a fast scanning speed can lead to varying degrees of porosity and unmelted powder being left inside the formed parts. From the temperature cloud maps of the cross-sections in Figure 6b and Figure 7b, we can see that under the conditions of a low energy density caused by the above two factors, the depth of the molten pool is 30 μm and 35.6 μm, respectively. Under the experimental conditions with a layer thickness of 40 μm, a good interlayer failed to form at a good overlap rate, and unmelted and sticky powder were present. The final formed parts showed varying degrees of small pores and cracks. With the increase in input laser energy density, the shape of the molten pool improved, and the density and microhardness gradually increased. With a further increase in laser energy density, the residual stress caused by large temperature gradients in the formed parts due to excessive energy density led to small cracks, and the density and microhardness of the formed parts decreased. In addition, under different laser energy densities, the density and microhardness of the formed parts showed the same trends.
In order to further investigate the effect of energy density on molding densities and microhardness, the thermal behavior of the melt pool needs to be studied. The degree of densification of molded parts is related to the liquid phase viscosity, liquid phase surface tension, and capillary action of the material during the molding process [29]. When the laser energy density is low, the melt pool temperature is low, the liquid phase viscosity is high, and the liquid pool fluidity is poor, resulting in a sticky powder with an insufficient lap ratio for the melt pool, which in turn produces large irregular pores, and due to the difference in thermal conductivity between the WC ceramic particles and the IN718 metal powders, the thermal conductivity of the melt pool is rapid in the vicinity of the WC particles, which leads to a reduction in the local temperature and an increase in viscosity in the melt pool compared with that of the liquid phase of the pure IN718 alloy. In addition, under the capillary effect caused by surface tension and Marangoni convection, there is a weak convection from the edge to the inside of the melt pool, and under the influence of this convection, the inside of the melt pool is unstable and prone to the formation of pores. In addition, when the laser energy density is low, the wettability of the WC ceramic particles with the metal substrate is low, resulting in poor bonding between the substrate and the reinforcing phase and low microhardness.
With the increase in laser energy density, the melt pool temperature is enhanced, the liquid phase viscosity is reduced, the surface tension and the liquid phase convection inside the melt pool are enhanced, the liquid phase metal expands uniformly, the morphology of the melt pool is stable, the densification and the overlap rate are increased, the metallurgical bonding between the ceramic-reinforcing phase and the metal substrate is enhanced, and the microhardness is enhanced. When the laser power is maintained at 150 W and the scanning speed is 250 mm/s, we can see from the temperature cloud map of the cross-section shown in Figure 8b that the depth and width of the melt pool are 42.12 μm and 125.18 μm, respectively. At this time, both the overlap rate of the melt and the overlap rate between layers meet the requirements, and the quality of the resulting product is the best, as shown in the metallographic analysis of Figure 8c, with no obvious pores or cracks.

3.2. Microstructure Analysis of Composite Materials

Figure 9 shows a microstructure image of the IN718/WC composite under a 150 W, 250 mm/s molding process. Figure 9a shows the WC-reinforced phase and its surrounding microstructure; the gray-white region in the center of Figure 9a is the WC particles, where a gradient transition layer was found to have formed between the WC particles and the IN718 matrix, with an average width of the transition layer of 4.51 μm, and the reinforced phases, such as gray-white γ′ and γ″, are uniformly dispersed among the gray-black metallic matrix γ-phase in the other regions [30,31,32], with an average size of up to 1 μm. The white phase A shown in Figure 9b was determined to have a complex (Ni, Cr, M)C carbide formed by Ni, Cr, Fe, and C as its main metal composition. Additionally, the sizes of the γ′- and γ″-reinforced phases were found to increase with the increase in WC content. Gu et al. [33] found that the carbon element in the metal-carbide ceramic particles is more likely to form metal carbides with Ni, Mo, and other elements during the molding process, which leads to diffusion of the elements and generates more precipitated phases, especially the reinforced phases of the metal carbides.
Figure 10 shows the elemental distribution of different thicknesses of the transition layer of WC particles, and the four points A, B, C, and D correspond to the WC-reinforced phase, the 1/3-thickness transition layer, the 2/3-thickness transition layer, and the boundary of the metal matrix, whose elemental profiles correspond to Figure 10a–d. From point A to point B, it can be clearly observed that the Cr, Ni, and Fe elements appear in the elemental maps, and from point (b) to point (d), the distance increases, and the peak of the W element disappears, but the peak of the C element still exists. At point A, the percentages of W and C are 91.95 wt% and 8.05 wt%, respectively. At point B, the percentages of W, C, Cr, Ni, and Fe are 65.19 wt%, 6.34 wt%, 11.89 wt%, 11.9 wt%, and 4.67 wt%, respectively. At point C, the percentages of C, Cr, Ni, and Fe are 3.1 wt%, 31.47 wt%, 47.26 wt%, and 18.17 wt%. The percentages of C, Cr, Ni, and Fe at point D are 2.34 wt%, 29.5 wt%, 49.35 wt%, and 18.8 wt%, respectively. We can see that with the increase in distance, the amount of W and C elements continues to decrease, in which the W element becomes seriously attenuated with distance, with nearly complete attenuation at 2/3 thickness, and the proportion of carbon element from the center of the WC to the transition layer and to the edge of the metal substrate is attenuated by 70.93%, which proves that the W and C elements in the WC during the molding process diffuse into the metal substrate, with more C elements diffusing into the metal substrate than W elements. In addition, the creation of the transition layer between the WC ceramic particles and the metal matrix also confirms that the reaction between WC and the elements in the alloy during the molding process leads to the diffusion of the elements between the reinforcing particles and the alloy matrix, which indirectly proves that the WC and the metal matrix form a metallurgical bond, leading to an improvement in the molding performance.
In order to further determine the phase composition of IN718/WC composites molded under the SLM process, the specimens doped with 1% WC were subjected to XRD analyses, and the patterns are shown in Figure 11.
From the XRD spectra of (a) pure IN718 alloy, obvious {111}, {200}, and {220} crystal diffraction peaks can be observed, which are analyzed as the Ni-γ phase; in addition, the γ′ phase with the chemical formula Ni3 (Al, Ti, Nb) and the γ″ phase with the chemical formula Ni3Nb, which were formed by Ni, Al, Ti, and other elements, were also identified, which coincides with the microstructural analysis of the material in a previous paper. In Figure 11b, the crystal diffraction peaks of {222}, {400}, {440}, and {602} can be observed, which are determined to be the metal carbides formed by Cr, Ni, Fe, and other metal elements with C, and the solid solutions of chemical formula NixWy with crystal diffraction peaks of {111}, {200}, {220}, and {311} are observed. This proves that the addition of WC ceramic particles in the molding process has a chemical reaction with the metal matrix, forming a metallurgical bond and exhibiting an elemental diffusion phenomenon, which is also consistent with the previous analysis. As the content of WC continues to increase, the XRD spectra show obvious WC phases with diffraction peaks of {111}, {200}, {220}, and {311}, which proves that with the increase in WC content, due to the difference in thermophysical parameters between WC and IN718, the temperature distribution inside the molten pool in the molding process is uneven, and the molten pool is not as stable as the molten pool of pure IN718 alloy, which leads to the formation of the molten pool at the same laser energy density, which is not as stable as the molten pool of pure IN718. As a result, under the same laser energy density, part of the powder cannot be fully melted and does not experience good metallurgical bonding with WC, resulting in an increase in the content of the WC phase, while the metallurgical bonding during the formation of the hard phase changes is not obvious.

4. Discussion

4.1. Effect of Molding Process on Friction and Wear Properties of Composites

Among the many SLM laser process parameters, laser power and scanning speed have the greatest influence on molding; these two parameters directly affect the energy input of molding and affect the temperature and duration of the liquid phase of the metal; and ultimately, the molding quality of the SLM-molded parts will be affected, so focusing on the study of the influence of laser power and scanning speed on the molding quality becomes necessary. Figure 12 shows the change curve of the friction coefficient of SLM-molded parts under different laser powers, from which we can see that the friction coefficient shows a tendency to first rise and then level off over time. The friction coefficient rises continuously at the early stage of molded part wear, at which time the friction process between the molded part and silicon nitride is in the stage of friction, and the friction coefficient of the specimen gradually levels off and fluctuates in small ranges as the duration under friction prolongs. With this prolonged friction time, the friction coefficient of the specimen gradually stabilizes and fluctuates within a small range. The time taken for the friction coefficient to stabilize at the early stage of wear and the fluctuation of the friction coefficient at the rising stage are related to the size and number of macroscopic pores, cracks, and other defects on the friction surface of molded parts. To ensure the accuracy of the experimental results, we conducted friction tests on five molded parts under the same process parameters. After the friction test entered the stable wear period, we plotted the average values of the samples at each time point and took the average friction coefficient during the stable wear period to create a bar chart.
Figure 12 shows that under three different laser powers, when the laser power is 150 W, the friction coefficient first rises to a stable value compared with the other laser power parameters and then the time taken for this stabilization decreases when the laser power rises from 75 W to 150 W, which is mainly due to the fact that internal defects in the molded parts improve with the increase in the laser power, and the number of generated cemented carbide phases increases with the increase in laser power. With the increase in laser power, the melting of the powder is sufficient, the metallurgical bonding of the powder is improved, and the number of generated cemented carbide phases increases; and when the laser power increases from 150 W to 225 W, the friction coefficient tends to increase stably, which is mainly due to the increase in laser power, resulting in the appearance of microscopic local cracks and spherical pores in the molded parts, and the decrease in densification.
The average coefficient of friction and wear of SLM-molded parts at different laser powers are shown in Figure 13.
From Figure 13, we can see that when the laser power is 150 W, the wear volume and the average friction coefficient are the smallest compared with other laser powers, 0.0371 mm3 and 0.61, respectively; in contrast, the average friction coefficient when the laser power is 75 W is 0.68, and the wear volume is 0.0393 mm3. With the increase in laser power from 75 W to 150 W, the average friction coefficient decreases by about 10.3%, and the wear volume decreases by about 5.6%. Along with the increase in laser power, the macroscopic defects inside the molded parts obviously improve, the densities increase, the microhardness increases, the ability of the surface of the molded parts to resist the extrusion effect of silicon nitride on the abrasive ball improves, the irreversible deformation under the continuous loading of the ball improves, the material loss of the molded parts in the friction experiments reduces, and the volume of the wear decreases. When the laser power increases from 150 W to 225 W, the average coefficient of friction increases by about 16.4% and the wear volume increases by about 9.7%. Along with the increase in laser power, the porosity of the molded parts increases, the microhardness decreases, the generation of abrasive chips accelerates under continuous loading of the hard grinding ball, the average coefficient of friction increases, and the wear volume increases.
In order to further investigate the mechanism of SLM laser power on the wear performance of molded parts, the following section will begin by presenting the wear morphology to study the relationship between laser power, the width and depth of the abrasion mark, and abrasion mark morphology. Research has found that the cross-sectional area of wear marks is generally similar to the wear mark length. The product of the cross-sectional area of wear marks after removing the head and tail and the wear mark length is defined as the wear amount. As shown in Figure 14, the highest point on both sides of the wear profile curve is connected horizontally, and a line segment is made perpendicular to the horizontal line, so that one end of the segment is the lowest point of the curve and the other end of the segment is on the horizontal line; then, the length of the horizontal line segment is the width of the abrasion mark, and the vertical line segment is the depth of the abrasion mark. From Figure 14, when the laser power is 75 W, 150 W, and 225 W, the widths of the abrasion mark are 0.9936 mm, 0.9061 mm, and 1.0306 mm and the depths of the abrasion mark are 19.023 μm, 18.985 μm, and 19.105 μm, respectively. With the increase in laser power, the abrasion mark depth and width show a tendency to first decrease and then increase, and the abrasion mark width and depth are the lowest when 150 W is used. The width and depth of the abrasion marks were the lowest at 150 W. When the power increases from 75 W to 150 W, the depth and width of the abrasion marks decrease by about 0.2% and 8.8%, respectively, and the abrasion mark morphology tends to be centrosymmetric and regular with the increase in laser power. When the laser power increases from 150 W to 225 W, the depth and width of the abrasion marks increase by about 0.6% and 12.1%, and the abrasion mark shape tends to be irregular with the increase in laser power. In addition, the sensitivity of abrasion depth to laser power is higher than that of abrasion width. This sensitivity is mainly due to the close relationship between the density and microhardness of the SLM-molded parts, with the density and microhardness reaching their maxima when the laser power reaches 150 W, which makes the molded parts more resistant to deformation under continuous loading of the grinding ball and attenuates the effect of the abrasive debris generated during the wear process on the abraded contact surfaces. At a laser power of 150 W, the abrasive and furrow wear reduce compared to at 75 W and 225 W, and the wear resistance of the specimens improves.
Figure 15 shows the change curves for the friction coefficients of the molded parts with time under different scanning speeds, from which we can see that at the early stage of wear, the molded parts are at the stage of friction from the small grinding balls; the contact area between the small balls and the molded parts is small, and the friction coefficients are irregular and gradually increase. With the prolongation of wear time, the friction coefficient of each molded part gradually tends toward the average friction coefficient, but in the late stage of the wear experiment, the friction coefficient of the molded part with low scanning speed fluctuates less, which is mainly due to the fact that with continuous wear, the WC particles in the molded part as well as other hard particles gradually become exposed and come into contact with the counter-abrasive grinding ball, which reduces direct contact with the substrate. In addition, the hard phases share some of this pressure, thus smoothing out the wear.
Figure 16 shows the average friction coefficient and wear amount of SLM-molded parts under different scanning speeds. We can see from Figure 16 that the average friction coefficient and wear amount of the molded parts become larger with a continuous increase in scanning speed. When the scanning speed is increased to 500 m/s, the average coefficient of friction is 0.62 and the wear amount is 0.0383 mm3, increases of about 1.6% and 3.1%, respectively, compared with those at 250 mm/s. When the speed is increased from 750 mm/s to 1000 mm/s, the average coefficient of friction and the wear amount of the molded parts are increased by 4.3% and 10.9%, respectively, which is a big difference between the average coefficient of friction and the wear amount of the molded parts at different speeds. The main reason for this difference is due to the density and microhardness of the molded parts. As mentioned in the previous section, the increase in laser scanning speed reduces the number of formed cemented carbide phases, and the wettability of the WC ceramic-reinforced phase with the metal substrate is reduced, meaning that a better metallurgical bond with the metal substrate cannot be formed, resulting in a reduction in microhardness and a gradual increase in the average friction coefficient and wear volume with scanning speed under the combined influence of these factors. In addition, the change in the average friction coefficient and wear volume with the scanning speed is consistent with that in densification and microhardness, which supports the conclusion that the average friction coefficient and the wear volume have a close relationship with densification and microhardness.
Figure 16 shows the changes in depth and width of the abrasion at different scanning speeds; from Figure 16, we can see that at scanning speeds of 250 mm/s, 500 mm/s, 750 mm/s, and 1000 mm/s, the widths of the abrasion marks are 0.9061 mm, 0.9245 mm, 0.9613 mm, and 1.0349 mm, respectively, and the depths are 18.985 μm. 19.017 μm, 19.017 μm, 19.017 μm, 19.017 μm, 19.017 μm, 19.017 μm, 19.092 μm, and 19.237 μm, respectively. With each increase in scanning speed, the change in abrasion width and depth also increases gradually: starting from 096061 mm and 19.985 μm, the width and depth of the abrasion increased by about 1.99% and 0.17% from a scanning speed of 250 mm/s to 500 mm/s; in contrast, the width and depth increased by 7.11% and 0.75%, respectively, when the scanning speed increased from 750 mm/s to 1000 mm/s, mainly due to the reduction in energy absorbed per powdered area with the increase in scanning speed, the metallurgy of the molten pool being poor, the porosity gradually rising, and the densification and microhardness gradually decreasing, resulting in a reduction in wear performance. The change in the width and depth of the abrasion marks with scanning speed is consistent with the change in density and microhardness with scanning speed. As the scanning speed increases, the shape of the abrasion marks gradually becomes irregular and asymmetric. This change in shape is mainly due to the scanning speed increasing, the wear from plastic deformation of the molding parts increasing, the size and number of internal defects increasing, the uneven pores during continuous loading of the grinding ball gradually rupturing and forming flaking debris, the wear surface flatness rapidly reducing, and the shape of the abrasion marks ultimately worsening.

4.2. Mechanism Behind the Effect of WC Content on the Wear Resistance of Molded Parts

The quality of IN718/WC composites molded following the SLM process is related to the molding process as well as the content of the ceramic-reinforcing phase added to the composite. Under the continuous action of a laser, the powder forms a molten pool, and the fluidity and convection inside the molten pool promote uniform distribution of WC particles in the molten pool, and with an increase in the content of WC in the ceramic-reinforcing phase; the opposite effect occurs on the fluidity and convection of the molten pool, which affects the molding quality. In order to determine the effect of WC content on the friction properties of the composites, nickel matrix composites were prepared by setting different mass fractions of WC powders during the molding process, with a laser power of 150 W and a scanning speed of 250 mm/s. The results showed that the WC powder was used in the molding process for the preparation of the nickel matrix composites.
The friction experiments were carried out on SLM-molded parts with different WC contents, and the friction coefficients under the dry friction experimental conditions are shown in the following graphs; from Figure 17, we can see that at the early stage of the friction experiments, the molded parts and the silicon nitride counter-abrasive ball are in the friction stage, and so, the friction coefficients show a fluctuating upward tendency, among which the fluctuation of the 10% WC/IN718 composites is more obvious, which is mainly related to the densities of the molded parts. The 10% WC/IN718 composite material has a higher porosity than other materials; the number and size of pore defects are larger; and the uneven morphology caused by the porosity easily forms abrasive debris during the dry grinding process, which makes the friction coefficient fluctuate more. With the friction experiment, the friction coefficient tends to stabilize, and the molded parts and the counter-abrasive ball enter into a stable friction period.
Figure 18 shows the average coefficient of friction and wear volume of SLM-molded parts with different WC contents, from which we can see that the average coefficient of friction and wear volume of the nickel matrix composites with WC additions are lower compared with that of the pure IN718 alloys, and are especially the lowest when the WC content is 5%. The average coefficient of friction at this time is 0.785 and the wear volume is 0.02973 mm3, which show that the WC, as a ceramic-reinforcing phase, can effectively reduce the friction coefficient and wear volume of the molded parts. This shows that WC, as a ceramic-reinforcing phase, can effectively reduce the coefficient of friction and wear volume of molded parts.
In addition, we can conclude from Figure 18 that when the WC content increases from 0% to 5%, the average friction coefficient and wear volume show a decreasing trend, decreasing by 26.95% and 4.27%, respectively, for a WC content of 5% compared with that of the pure IN718 alloy. This decrease is mainly due to the excellent properties of WC, which help lubricate and reduce wear during friction and the wear between the silicon nitride on the grinding ball and the molded parts. With an increase in WC content, the average coefficient of friction and wear volume of molded parts showed a rising trend; compared with 5% WC content, the average coefficient of friction and wear volume of molded parts increased by 13.57% and 2.12% when WC content was 10%, which was mainly due to the fact that with the increase in WC content, the densification of the molded parts decreased; the number of pores and the size of the porous space increased; and in the case of silicon nitride subjected to abrasive wear, the average friction coefficient and wear volume increased by 4.27%. The main reason for these changes is that with the increase in WC content; the density of the molded parts decreases; the number and size of pores increase; the concave–convex morphology on the contact surface of the silicon nitride grinding balls and molded parts increases; and the number of instances of flaking in the molded parts increases during abrasion under continuous pressure from the grinding balls, which accelerates the abrasion.
In order to further study the influence of WC content on the friction and wear performance of SLM-molded parts, the depth and width of wear marks were studied together with the wear morphology. From Figure 19, we can conclude that the variation rules of abrasion depth and width of molded parts with WC content are more or less the same, and this trend is the same as that of microhardness with WC content. When the content of WC increases from 0% to 5%, the depth and width of the abrasion mark decreases significantly, which is consistent with the change in wear amount, thus reflecting that the addition of WC has a significant effect on the improvement in wear performance. This study focuses on the wear resistance level of materials during stable wear processes. Through a statistical analysis of multiple samples, we can determine the effects of laser power, scanning speed, and WC content on wear resistance. In terms of WC content, the minimum difference in the average wear resistance coefficient between WC samples with different contents is about 5%. Although the difference is small, the results can summarize the trend for the influence of WC content on the wear resistance of formed parts. When the WC content is increased from 5% to 10%, the wear increases, flaking of the molded parts increases during the wear process, and the high temperature generated by the continuous pressure and back-and-forth friction of the silicon nitride on the grinding balls causes the abrasive chips to adhere to the surface of the abrasive marks, which leads to increased wear and, ultimately, an increase in the depth and width of abrasive marks.
As can be seen from Figure 19, the wear pattern after changing the WC content is more regular and symmetrical than that after changing the SLM laser parameter, and the wear pattern after increasing the WC content from 0% to 5% is more regular than that after increasing the content from 5% to 10%, which is mainly due to the fact that with the increase in WC content, the densification decreases, and the number and size of pores in the molded parts increase, leading to an increase in wear heat and increases in the depth and width of the wear marks. The main reason for these changes is that as the WC content increases, the densification decreases and the number and size of pores in the molded part increase, which leads to an increase in wear heat generation and an increase in the adhesion of abrasive debris and ultimately affects the regularity of the wear pattern.

5. Conclusions

In this paper, IN718/WC composite specimens were successfully prepared using SLM additive manufacturing technology. Firstly, the effect of different laser powers on the molding quality was analyzed via numerical simulation, and the appropriate laser power interval was determined for the experiment. The effects of different laser powers and scanning speeds on the microstructure, microhardness, and friction properties of IN718/WC composites were investigated, and the mechanism of WC enhancement on the wear resistance of IN718 was analyzed. The main conclusions of this study are as follows:
(1)
We found that the best molding quality of the composites was achieved when the laser power was 150 W and the scanning speed was 250 mm/s. Under this parameter, the liquid phase existed for a moderate period of time, the degree of powder melting was optimal, and the overlap rate between the melting channels and the adjacent powder layers was better, so that a high-quality metallurgical bond could be formed.
(2)
Through a microstructure analysis of the composite material, the microstructure of the composite material showed that the γ′ and γ″ reinforcing phases are dispersed in the Ni–γ matrix phase, the average size of the reinforcing phases is about 1 micron, WC ceramic particles are uniformly distributed in the matrix phase, the WC and metal matrix produce in situ reactions and form a transition layer, and the elements with a gradient distribution of the transition layer can effectively improve the microhardness and wear resistance of the molded parts.
(3)
When the laser energy density is relatively low, the input energy of the laser heat source is low, the viscosity of the liquid phase is high, and the liquid phase convection inside the melt pool is poor, which reduces the capillary force on the WC ceramic particles and exacerbates the agglomeration of the nanoparticles, resulting in poor quality in the formed specimens.
(4)
With the increase in WC content, the molding densities of the IN718/WC composites decreased and their microhardness increased. When the WC content was 5%, the molded parts showed optimal wear resistance, the friction coefficient fluctuated steadily during the friction process, the degree of wear was low, and the wear volume was reduced to 0.02973 mm3. The average friction coefficient and the wear volume of the molded parts decreased by 26.95% and 4.27%, respectively, compared with that of the pure IN718-molded parts with the same WC content.
In this study, the molding process parameters of IN718/WC composites were investigated; the microstructure, microhardness, and friction and wear properties of the composites were systematically analyzed under different laser powers and scanning speeds; and the mechanism behind the enhancement of the WC-enhanced phase was analyzed, which will provide experimental references and theoretical guidance for improvements to the mechanical properties of nickel-based high-temperature alloys in the future.

Author Contributions

Conceptualization, Y.S. and X.Z. (Xuejin Zhao); methodology, Y.X.; software, X.Z. (Xiaoyu Zhao); validation, X.Z. (Xiaoyu Zhao) and X.Z. (Xuejin Zhao); formal analysis, X.Z. (Xiaoyu Zhao); investigation, Y.X.; resources, X.Z. (Xiaoyu Zhao); data curation, Y.X.; writing—original draft preparation, X.Z. (Xiaoyu Zhao); writing—review and editing, Y.X. and X.Z. (Xuejin Zhao); visualization, X.Z. (Xiaoyu Zhao); supervision, X.Z. (Xuejin Zhao); project administration, Y.S. and X.Z. (Xuejin Zhao); funding acquisition, Y.S. All authors have read and agreed to the published version of the manuscript.

Funding

This research was financially supported by the Shandong Province Natural Science Foundation (grant ZR2020ME162).

Data Availability Statement

The original contributions presented in the study are included in the article, further inquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

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Figure 1. XRD spectra of the IN718 alloy.
Figure 1. XRD spectra of the IN718 alloy.
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Figure 2. (a) IN718 powder morphology; (b) WC powder morphology.
Figure 2. (a) IN718 powder morphology; (b) WC powder morphology.
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Figure 3. Huashu Hi-Tech FS271M constituent selective laser melting equipment.
Figure 3. Huashu Hi-Tech FS271M constituent selective laser melting equipment.
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Figure 4. Scanning strategy schematic.
Figure 4. Scanning strategy schematic.
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Figure 5. (a) Variation in densification with laser energy density; (b) variation in microhardness with laser energy density.
Figure 5. (a) Variation in densification with laser energy density; (b) variation in microhardness with laser energy density.
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Figure 6. (a) Top view of a temperature field simulation with power at 100 W and scanning speed at 250 mm/s; (b) cross-sectional view; (c) microstructure.
Figure 6. (a) Top view of a temperature field simulation with power at 100 W and scanning speed at 250 mm/s; (b) cross-sectional view; (c) microstructure.
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Figure 7. (a) Top view of a temperature field simulation with power at 150 W and scanning speed at 350 mm/s; (b) cross-sectional view; (c) microstructure.
Figure 7. (a) Top view of a temperature field simulation with power at 150 W and scanning speed at 350 mm/s; (b) cross-sectional view; (c) microstructure.
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Figure 8. (a) Top view of a temperature field simulation with power at 150 W and scanning speed at 250 mm/s; (b) cross-sectional view; (c) microstructure.
Figure 8. (a) Top view of a temperature field simulation with power at 150 W and scanning speed at 250 mm/s; (b) cross-sectional view; (c) microstructure.
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Figure 9. (a) Microstructure image of SLM-molded IN718/WC composite material; (b) phase A magnified view.
Figure 9. (a) Microstructure image of SLM-molded IN718/WC composite material; (b) phase A magnified view.
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Figure 10. Microstructure calibration point elemental mapping of IN718/WC composites: (a) point A mapping; (b) point B mapping; (c) point C mapping; and (d) point D mapping.
Figure 10. Microstructure calibration point elemental mapping of IN718/WC composites: (a) point A mapping; (b) point B mapping; (c) point C mapping; and (d) point D mapping.
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Figure 11. (a) XRD spectra of IN718; (b) XRD spectra of composites doped with 1% WC.
Figure 11. (a) XRD spectra of IN718; (b) XRD spectra of composites doped with 1% WC.
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Figure 12. Friction coefficient versus time curves at different laser powers.
Figure 12. Friction coefficient versus time curves at different laser powers.
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Figure 13. (a) Average coefficient of friction at different laser powers. (b) Wear of molded parts at different laser powers.
Figure 13. (a) Average coefficient of friction at different laser powers. (b) Wear of molded parts at different laser powers.
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Figure 14. Cross-sectional view of abrasion mark morphology at different laser powers.
Figure 14. Cross-sectional view of abrasion mark morphology at different laser powers.
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Figure 15. Variation in the friction coefficient of molded parts over time at different scanning speeds.
Figure 15. Variation in the friction coefficient of molded parts over time at different scanning speeds.
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Figure 16. Cross-section of abrasion marks at different scanning speeds.
Figure 16. Cross-section of abrasion marks at different scanning speeds.
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Figure 17. Friction coefficient curve of molded parts with different WC contents.
Figure 17. Friction coefficient curve of molded parts with different WC contents.
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Figure 18. (a) Average coefficient of friction of SLM parts with different WC contents. (b) Wear of SLM parts with different WC contents.
Figure 18. (a) Average coefficient of friction of SLM parts with different WC contents. (b) Wear of SLM parts with different WC contents.
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Figure 19. Wear profiles of molded parts with different WC contents.
Figure 19. Wear profiles of molded parts with different WC contents.
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Table 1. Main chemical composition of IN718.
Table 1. Main chemical composition of IN718.
ElementContent (%)
Nb4.75~5.5
Ti0.65~1.15
Mo2.8~3.3
C≤0.08
Cr17~21
Ni50~55
Co≤1.0
Al0.2~0.8
FeResidual
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Zhao, X.; Zhao, X.; Xu, Y.; Shi, Y. Friction Performance Analysis of WC-Reinforced IN718 Composite Material Based on SLM Process. Metals 2024, 14, 1361. https://doi.org/10.3390/met14121361

AMA Style

Zhao X, Zhao X, Xu Y, Shi Y. Friction Performance Analysis of WC-Reinforced IN718 Composite Material Based on SLM Process. Metals. 2024; 14(12):1361. https://doi.org/10.3390/met14121361

Chicago/Turabian Style

Zhao, Xuejin, Xiaoyu Zhao, Youfan Xu, and Yongjun Shi. 2024. "Friction Performance Analysis of WC-Reinforced IN718 Composite Material Based on SLM Process" Metals 14, no. 12: 1361. https://doi.org/10.3390/met14121361

APA Style

Zhao, X., Zhao, X., Xu, Y., & Shi, Y. (2024). Friction Performance Analysis of WC-Reinforced IN718 Composite Material Based on SLM Process. Metals, 14(12), 1361. https://doi.org/10.3390/met14121361

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