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Article

Influence of Powder Mass Flow Rates on Wear Resistance and Impact Toughness of Inconel 718 Surface Coatings

by
Chaoming Zhang
1,
Linsen Shu
1,*,
Peiyou Li
2,
Jiangtao Gong
1,
Wei He
1 and
Helong Yu
3
1
School of Mechanical Engineering, Shaanxi University of Technology, Hanzhong 723001, China
2
School of Materials Science and Engineering, Shaanxi University of Technology, Hanzhong 723001, China
3
National Key Laboratory for Remanufacturing, Army Academy of Armored Forces, Beijing 100072, China
*
Author to whom correspondence should be addressed.
Coatings 2023, 13(11), 1877; https://doi.org/10.3390/coatings13111877
Submission received: 7 September 2023 / Revised: 1 October 2023 / Accepted: 5 October 2023 / Published: 1 November 2023
Browse Figures
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Abstract

:
Optimum laser cladding processing parameters were obtained based on the study of various powder mass flow rates for the repair of 27SiMn steel parts using Inconel 718 powder. In this study, fusion coating process parameters were set according to the influence of powder mass flow rates on fusion coating properties. The cross-sectional microstructure, X-ray diffraction patterns, micro-hardness distribution, friction, wear properties, impact properties, and fracture morphology of the clad layers were investigated for the various process parameters. The results indicate that the volume fraction for the Ni3Fe phase increases and then decreases as the powder mass flow rate increases. The micro-hardness, friction properties, and impact properties of the samples followed the same trend as those of Ni3Fe. At powder mass flow rates up to 20 g/min, microstructures in the cladding layer tended to be uniform with microstructures with fine grains, and micro-hardness and impact toughness reached maximum values of 328.2 HV0.5 and 45.4 J/cm2, respectively, which show better mechanical properties and wear resistance.

1. Introduction

Because of its excellent comprehensive properties, 27SiMn steel is widely used in petrochemical equipment [1], shipbuilding [2], machining, and other applications of high-strength engineering [3]. These 27SiMn steel parts work in high-strength environments such as those involving severe wear and impact for a long time, and their surface structure becomes damaged, which makes such parts fail or unable to work normally in serious cases, thus bringing significant economic losses [4,5]. Marushchak et al. [6] investigated the impact toughness of monometallic and bimetallic specimens cut on the rolls of a continuous casting machine with 35G2 steel fused to 18Kh11MNFB steel. The results showed that the energy required for crack initiation in fused 18Kh11MNFB steel was higher than that of 35G2 steel in the roll body itself [6]. Similarly, using laser cladding technology to repair worn surfaces or strengthen existing metal surfaces to obtain wear-resistant and impact-resistant materials can significantly improve the service performance and longevity of these components [7,8].
Currently, Fe-based [9], Ni-based [10], and Co-based [11] alloy coatings are often used as protective coatings on metal surfaces. In particular, they have a high intensity and hardness, wide application range, and a low cost is associated with Fe-based laser cladding coatings, so they have attracted much attention [12]. However, due to the melting point and low fluidity, high-hardness Fe-based coatings are susceptible to cracking defects during the rapid cooling and rapid heating processes of laser cladding, resulting in a low strength of the cladding layer, which affects its popularization and application [13,14]. Surface damage on 27SiMn steel parts can be repaired using Fe-based powder cladding, but the surface strength after repair is low [15] and cannot meet parts’ service requirements. Thus, to ensure the quality of repairs, it is necessary to select a high-performance powder–material pair. Inconel 718 alloy offers a good combination of properties, including outstanding strength and toughness, resistance to oxidation, creep resistance, and high fatigue life [16]. In particular, the elements Ni and Nb provide high solid solution strengthening of the alloy, which enhances its corrosion resistance, strength, and toughness [17,18] and is particularly applicable to the enhancement and repair of 27SiMn steel part surfaces. The properties of laser cladding coatings are also largely influenced by the parameters of the laser cladding process. Xie et al. [19] investigated the influence of different laser powers on a Fe314 alloy cladding layer’s microstructure and its properties. The results showed that the grain size of the coating gradually increased and the micro-hardness gradually decreased with increasing laser power [19]. Dariusz et al. [20] investigated the influence of different powder mass flow rates and laser powers on the properties of a Fe/WC metal matrix composite coating. The results showed that increasing the powder mass flow rate and the laser power can increase coating thickness and that increasing the powder mass flow rate appropriately can also improve the surface hardness of the coating [20]. Kim et al. [21] investigated the effects of powder mass flow rate and laser power on the mechanical properties and interfacial cracking of 630 stainless steel coatings, finding that decreasing the powder mass flow rate moderately was more effective in suppressing coating defects than changing the laser power, and the tensile strength of the cladding samples whose interfacial defects were eliminated by lowering the powder mass flow rate was as high as 97% of that of the substrate [21]. Appropriate process parameters are essential to ensure excellent coating performance. There have been many studies on the effects of process parameters on the performance of coatings; however, after repair, the strength of coatings has not been not satisfactory [22]. At present, there have been few studies on the influence of powder mass flow rates on the coating properties of Inconel 718 alloy, which hinders its engineering application in laser cladding remanufacturing. Therefore, in this study, optimum process parameters for Inconel718 alloy laser cladding coatings were explored to understand the effects of powder mass flow rates on the properties of fused Inconel 718 coatings. Different powder mass flow rates were used for the preparation of Inconel 718 coatings on 27SiMn steel surfaces, keeping other process parameters unchanged. The cross section microstructure, XRD diffraction pattern, micro-hardness distribution, friction and wear properties, impact properties, and fracture morphology of the coating were analyzed at different powder mass flow rates. This study provides a valuable reference for engineering applications of 27SiMn steel parts’ laser cladding remanufacturing and Inconel 718 cladding layer structure and performance control.

2. Experimental Procedure

The base material was 27SiMn steel (chemical composition was as follows: 0.30 wt% Ni, 0.30 wt% Cr, 0.15 wt% Mo, 0.32 wt% C, 0.30 wt% Cu, 1.20 wt% Mn, 1.25 wt% Si, and the remainder comprised Fe), the surface of which was ground flat before the cladding test, and the oil and dust on the surface were removed with anhydrous ethanol. The cladding powder was Inconel 718 alloy powder (chemical composition was as follows: 0.51 wt% Ni, 0.19 wt% Cr, 0.5 wt% Nb, 0.03 wt% Mo, 0.01 wt% Ti, 0.00034 wt% C, 0.0005 wt% Cu, 0.0014 wt% Co, 0.0014 wt% Mn, 0.0004 wt% Si, and the remainder comprised Fe), with the particle size ranging between 45 and 106 μm.
The laser cladding experiments were carried out using ABB’s robotic laser cladding system, which uses a YLS-3000 semiconductor laser (wavelength 1064 nm) to provide the laser energy. The powder feed mode was “light powder coaxial”, and the powder gas and shielding gas were high-purity argon gas. In order to study the effect of the powder mass flow rate on the coating properties of the fused Inconel 718 cladding layer, fusion cladding experiments were carried out using the process parameters shown in Table 1. Two-layer multi-channel Inconel 718 coatings were produced with 50% overlap and “bow”-shaped paths on the 27SiMn steel surface (as shown in Figure 1).
The samples’ preparation process is shown in Figure 2, and cladding samples under each set of laser cladding process parameters are shown in Figure 2a. Samples were cut out using a wire-cut EDM machine according to the sampling position in the figure. The impact specimen’s size was 10 mm × 10 mm × 55 mm, the coating thickness was 1.5 mm, and the specimen had a V-shaped notch of 2 mm depth according to ASTM E23, as shown in Figure 2b. The specifications and dimensions of the metallographic and wear specimens are shown in Figure 2c and Figure 2d, respectively.
After the metallographic specimens were polished and etched (etching solution was a 3:1 solution of hydrochloric acid and HNO3), cross-sectional and microstructural features were observed and analyzed for metallographic specimens at different powder mass flow rates using a VHX-7000 microscope. X-ray diffraction analyses were then carried out. After the impact and friction wear samples were taken, the clad surface was polished and leveled using a universal tool grinder. The impact work for the substrate and cladding specimens was then determined using a Charpy pendulum impact tester, and the coefficient of friction curves for the substrate and cladding specimens were determined using an HT-1000 friction wear tester. Finally, the wear and impact fracture patterns were observed using a VHX-7000 microscope and a JSM-6390LV scanning electron microscope.

3. Results and Discussion

3.1. Microstructure of Cladding Layer Cross-Section

Figure 3a–e presents photographs of the cross-sectional morphology of cladding specimens prepared at powder mass flow rates of 16 g/min, 18 g/min, 20 g/min, 22 g/min, and 24 g/min. The cross-sectional morphology of the fused cladding layers was found to be largely similar across conditions of various powder mass flow rates; however, variations were observed in the width, height, and depth of the heat-affected zone for the fused cladding layers, as highlighted in Table 2. As the mass flow rate of the powder increases, the coating width and thickness gradually increase. Concurrently, the depth of the melt pool gradually decreases. The depth of the heat-affected zone experiences a reduction followed by an increase. This is due to the fact that the powder is conveyed at a slower speed, absorbing less energy, allowing most of the laser energy to be applied to the substrate and creating a greater depth of the melt pool with more heat accumulating in the bottom part of the pool. Consequently, the heat-affected zones become deeper. Increasing the amount of powder progressively widens and elevates the cladding, leading to a reduction in power exertion over the substrate, a shallower melt pool, decreased heat buildup at the bottom of the pool, and a smaller heat-affected zone. As the powder mass flow rate increases, the width and height of the molten coating also increase, while the amount of energy acting on the substrate decreases. This leads to a shallower melt pool, with most of the laser energy melting the powder and causing it to enter the pool. As a result, there is an accumulation of heat at the bottom of the melt pool, which leads to a larger heat-affected zone. Overall, at a powder mass flow rate of 20 g/min, the depth of the melt pool is moderate and the fusion coating is firmly joined to the substrate, resulting in a shallow heat-affected zone. Therefore, the substrate is less likely to be damaged by laser energy.
Figure 4 reveals the microstructure of the top, transition zone, and bottom of the fused cladding layer at different powder mass flow rates. The cladding exhibits a dense and well-formed microstructure with apparent growth characteristics. The liquid metal solidifies from the bottom to the top of the bath due to the melt pool’s thermal conductivity. During continuous laser scanning, the top of the first solidified cladding layer acts as a matrix, allowing for clear distinction between the boundary of the intermediate layer and the cladding layer through the characterization of the transition zone. Micro-cracks were observed at the bottom of the cladding layer when powder mass flow rates were 16 g/min and 18 g/min, as shown in Figure 4a,b. A low powder mass flow rate is responsible for the excessive energy applied to the first layer during the melting of the second layer using the laser, resulting in cracking of the bottom of the melted cladding due to excessive heat buildup [23].
The cladding’s microstructure is influenced significantly by the temperature gradient (G) and solidification rate (R). As there are varying temperature gradients at different cladding layer locations during cooling, the microstructure differs according to positional differences [24]. Larger temperature gradients (G) near the substrate result in faster cooling. The liquid metal grows epitaxially at the bottom of the melt pool, displaying typical columnar growth. However, the growth of unsolidified columnar crystals is impeded by solidified cellular crystals and dendritic crystals because of an uneven heat dissipation and uneven solidification rate (R) in the bottom of the cladding [25]. Consequently, a combination of coarse columnar and dendritic crystals with a few cellular crystals eventually form in the lower section of the cladding. This is demonstrated by Figure 4a3–e3. The largest grain sizes observed at the base of the S1–S5 cladding coatings were 196 μm, 117 μm, 109 μm, 163 μm, and 262 μm.
The transition zone’s structure is comparable to that of the bottom of the coating. During the multi-layer cladding process, the solidified upper part of the initial cladding acts as the substrate for the second cladding layer. The solidified grains in the initial layer can continue to grow at the base of the second layer, competing with the unsolidified grains in the same layer [26]. Eventually, in the transition zone, they crystallize to form a mixed structure with dendrites being the primary component, along with a few columnar and cellular crystals. This is demonstrated in Figure 4a2–e2, where the maximum grain sizes in the transition zone of S1–S5 cladding coatings are 114 μm, 112 μm, 82 μm, 124 μm, and 144 μm. In Figure 4a–e, the second layer of the S1–S5 cladding coatings demonstrates the presence of a few large-sized columnar crystals that extend from the transition zone to the top. This occurrence is due to the significant temperature gradient (G) and the low solidification rate (R), which allow the grains to grow adequately. Furthermore, the columnar grain sizes of the S1~S5 cladding coatings measure 383 μm, 329 μm, 192 μm, 564 μm, and 589 μm.
Under the influence of both the protective gas and the heat dissipation of the solidified alloy, the heat flow at the top of the cladding alters, leading to a reduction in the temperature gradient (G) and an increase in the rate of solidification (R). Many crystal nuclei in the middle grew freely and ultimately coalesced to form a combination of short dendrites and cellular crystals on the top of the cladding, as depicted in Figure 4a1–e1. S1–S5 cladding coatings have maximum grain sizes of 72 μm, 63 μm, 58 μm, 65 μm, and 81 μm. Consequently, when the powder mass flow rate rises, the particle size of each component of the molten coating decreases initially and then increases.
When the powder mass flow rate is low, more laser energy is applied to the bonding zone, resulting in excessive laser energy and a fully grown cladding layer bottom and some grains in the transition zone with large G/R values. However, as the powder mass flow rate increases to 20 g/min, the heat used to melt the powder increases, reducing the energy acting on the substrate which leads to saturation of the laser energy. At this stage, the ratios of solidification growth rate to nucleation rate are low, causing a decrease in crystal grain size and an increase in the number of dendritic and columnar grains at the bottom of the cladding and at the transition zone. Continuous increases in powder mass flow rate correspond to decreasing energy acting on the substrate and increased laser absorption in the surrounding powder space. As the powder moves from the solid state to the liquid melt pool, metal flow and diffusion decline [27], causing heat to accumulate and boosting the G/R value. Ultimately, sizeable continuous coarse dendritic and columnar crystals take shape at the bottom of the cladding and the transition region.

3.2. XRD Analysis

Figure 5 displays the XRD patterns of cladding coatings S1–S5, which were prepared using substrate S0 and powder mass flow rates of 16 g/min, 18 g/min, 20 g/min, 22 g/min, and 24 g/min. Diffraction peaks of the γ-(Ni, Fe, Cr) phase are evident on the (110), (200), and (211) crystal surfaces of substrate S0 and cladding samples S1–S5. At the (111), (102), and (220) crystal surfaces of cladding coatings S1–S5, diffraction peaks of the γ″-Ni3Fe phase are also visible. When laser energy is applied to the metal powder, some of the Fe and Ni in the liquid metal create γ″-(Ni, Fe) phase at 1440 °C. After cladding, when the temperature reduces to 347–517 °C, solid phase transformation takes place in the γ″-(Ni, Fe) phase, leading to the formation of Ni3Fe phase. [28].
In the cladding samples depicted in Figure 5, the height of the diffraction peaks for the γ-(Ni, Fe, Cr) phases remained comparatively stable as the powder mass flow rate increased. However, a significant change was observed in the height of the diffraction peaks for the Ni3Fe phase. The total integrated area of the phase in the XRD spectra was used to estimate the continuous phase volume fractions. The peak integration area for (Ni, Fe, Cr) and Ni3Fe phase diffraction was determined using a peak-fitting procedure with the Pearson Ⅶ function. Refs [29,30] and Equation (1) were utilized to estimate the volume fraction ( V N i 3 F e ) of the Ni3Fe phase [31]:
V N i 3 F e = S N i 3 F e S ( Ni , Fe , Cr ) + S N i 3 F e
where S ( Ni , Fe , Cr ) and S N i 3 F e represent the integrated areas of the diffraction peaks of the γ-(Ni, Fe, Cr) and Ni3Fe phases, respectively. According to our calculations, samples S1–S5 display a greater relative volume fraction of the Ni3Fe phase compared to substrate S0, specifically 1.05%, 1.18%, 15.02%, 11.68%, and 4.71%. It has been observed that during the laser cladding process, the formation of the Ni3Fe phase initially increases and then decreases as the powder mass flow rate in the samples rises.

3.3. Micro-Hardness

Figure 6 presents the micro-hardness distribution of cladding coatings S1 to S5 for five powder mass flow rates. The results indicate a consistent decrease in coating hardness from the surface towards the substrate with varying powder mass flow rates. The adhesion area between the substrate and coating displays stable coating hardness.
When the powder mass flow rate ranges from 16 g/min to 24 g/min, the coating’s highest surface hardness varies from 317.2 HV0.5 to 315.4 HV0.5. The results illustrate a pattern of first increasing and then decreasing hardness with an increase in powder mass flow rate. There was little change observed in the hardness distribution of both the matrix and the heat-affected zone with an increase in powder mass flow rate. On the other hand, it is generally understood that a metal’s strength has a negative correlation with grain size, meaning that the smaller the grain size, the greater the strength and, therefore, the hardness of the material [32]. Alternatively, increasing the powder mass flow rate improves the solid solution and deforms the lattice size during non-dynamic equilibrium laser cladding processes, resulting in increased coating hardness [33]. The best coating hardness and overall performance are achieved when the powder mass flow rate is set at 20 g/min.

3.4. Friction and Wear Property

The graphs in Figure 7a depict the time-dependent friction coefficient curves of cladding coatings S1–S5 and the substrate S0 with time under different powder mass flow rates. The corresponding comparison of wear loss is shown in Figure 7b. The average coefficient of friction and wear loss of substrate S0 were 0.75 and 0.182 g, respectively. When the powder mass flow rate was 16 g/min, 18 g/min, 20 g/min, 22 g/min, and 24 g/min, there were reductions of 12%, 15%, 31%, 9%, and 4%, respectively, in the average friction factor of the coating compared to that of the substrate. Similarly, the wear loss was decreased by 43%, 46%, 64%, 38%, and 30% respectively, when compared to the substrate. The findings demonstrate that coatings S1–S5 have lower average friction factors and wear losses than the substrate. These results suggest that the coatings show better wear resistance than the substrate. Additionally, the average coefficient of friction and wear loss of each coating initially increase and then decrease with an increase in powder mass flow rate. This indicates that the abrasion resistance of each coating increases at first but then declines as the powder mass flow rate increases. When the powder mass flow rate is 20 g/min, the resulting coating exhibits superior hardness and abrasion resistance.
Figure 8 illustrates the wear morphology of coatings S1–S5 at specific powder mass flow rates as well as that of substrate S0. The surface of the substrate S0 sample reveals several pits and wear chips, as depicted in Figure 8a. During wear, the specimen’s surface’s contact point generates high temperatures that instantly soften it, leading to a reduction in local metal strength, adhesion transfer, and laminar spalling under shear deformation force. This phenomenon results in the formation of more abrasive debris and pitting, causing severe wear and tear [34]. If the powder mass flow rate is too low or too high, the coating exhibits low micro-hardness and a high coefficient of friction. The surface of the coating displays a high degree of roughness, as evidenced by the presence of numerous pits and furrows. This is further reinforced by the illustrations provided in Figure 8b,c,e,f which indicate that the surface stress of the coating has exceeded its fatigue strength. During the process of wear loss, the wear chips that are shed act as hardware particles, leading to the formation of surface cracks and eventually resulting in local fatigue, fracture, and spalling of the surface materials. This, in turn, contributes to the creation of craters which are indicative of particle wear, fatigue wear, and spalling wear mechanisms, respectively [35]. It can be inferred from these observations that the abrasion resistance of the coating is poor.
When the powder mass flow rate is 20 g/min, the coating exhibits greater micro-hardness and a lower friction coefficient. Initially, tiny fragments of metal on the coating’s surface flake off, producing abrasive particles that create deep grooves on both the coating and areas of the adhesion layer. Later on, some of the adhesion layer came off, and only a few adhesion pits emerged on the surface coating. The area of brittle spalling was limited, and abrasive wear was the primary mechanism. When compared to cladding coatings at different powder mass flow rates, the coatings with a powder mass flow rate of 20 g/min had shallow groove marks and fewer bonding layers, as portrayed in Figure 8d. It has been demonstrated that the coatings produced at this powder mass flow rate possess superior resistance to abrasion.

3.5. Impact Toughness

Figure 9 illustrates a comparison of impact toughness of cladding specimens (S1–S5) and substrate S0. The impact toughness of specimens S1 and S5 is lower than that of the substrate S0 whereas their impact resistance is weaker compared to it. Samples S2, S3, and S4 exhibit an increase in impact toughness of 15.6%, 39.3%, and 1.8%, respectively, in comparison to the substrate S0. The impact toughness (KCV) of S3 was the highest among all specimens, measuring at 45.4 J/cm2, and it showed the strongest impact resistance. This can be attributed to the formation of Ni3Fe phase during the laser melting of Inconel 718 alloy powder. The compound, created through the melting of Ni and Fe elements, enhances the thermodynamic stability of the coating and ultimately bolsters the strength of the cladding specimens [36]. At varying powder mass flow rates, specimen S3 displays a significantly higher relative volume fraction of Ni3Fe phase compared to the other cladding specimens, resulting in its superior strength. The impact resistance order of the cladding specimens is S3, S2, S4, S5, S1. The Hall–Petch effect [37] suggests that modifications to the grain size of the internal structure of the coating produce alterations in its impact properties. In the cladding, microstructural grains are refined, enhancing the impact resistance of the specimen. Nonetheless, larger microstructure particles in the cladding may render the specimen more prone to fracture when subjected to stress, particularly when the powder mass flow rate is too low or too high, resulting in diminished impact resistance. As the powder mass flow rate increases, the impact resistance of specimens rises and then declines.
Figure 10 depicts the three-dimensional fracture morphology of impact specimens S1–S5 at various powder mass flow rates, including a schematic diagram of typical zones and corresponding area histograms. As illustrated by Figure 10a–e, the fractures of specimens S1, S2, S4, and S5 appear relatively flat, with height differences only ranging between 1318 and 1720 μm. This suggests that the plastic work consumed during fracture by the specimens is minimal. Contrarily, the entirety of specimen S3 displays an uneven fracture, with a height differential of 12,181.32 μm. This suggests that specimen S3 absorbed considerable plastic work during fracture, leading to high levels of impact toughness that are consistent with the findings of the impact toughness analysis of cladding specimens S1–S5 in Figure 9.
The fracture process of impact specimens S1–S5 under identical impact loads is a continuous stage of shear deformation, crack inception, and propagation. The impact pendulum makes contact with the specimen and deforms it which correlates to the crack’s initiation point when deformation reaches its maximum level. Afterwards, the crack propagates, creating a shear lip and the final fracture zone of the specimen. These processes lead to the formation of corresponding regions on the fracture surface. As illustrated in Figure 10a1–e1, four distinct regions can be identified in the zones displaying characteristic patterns: (I) the shear region, (II) the shear lip region, (III) the subcritical crack growth region, and (IV) the final fracture region. It is important to note that these regions are clearly delineated and identifiable based on the patterns displayed in the figure.
The area histograms of the zones with typical fracture patterns, as illustrated in Figure 10a2–e2, reveal that the shear zone and shear lip zone of specimens S1–S5 are minimal. This indicates that only slight shear deformation transpired when the specimens were subjected to significant impacts. On the other hand, the shear zone and shear lip zone of specimen S3 are considerably more substantial than those of other specimens. This implies that specimen S3 underwent significant deformation before fracturing under impact load. However, the deformation of other specimens prior to crack initiation was negligible, resulting in poor impact resistance. Subsequently, the crack propagated from the central region of the specimen towards the ultimate fracture zone and shear lip zone. The crack growth zone in specimens S1–S5 was extensive, comprising 40% to 50% of the overall fracture area. The morphology, dimensions, and development of the crack growth zone were dependent upon the ability of the metal structure to resist shear deformation. The final fracture zone is the result of the low toughness fracture of the material, accounting for only 26.3% of the sample S3, while the final fracture zone of other samples is higher than S3.
Figure 11 illustrates the planar fracture morphology of impact specimens S1–S5 at various powder mass flow rates and their corresponding local enlargements. A closer examination reveals that the fractures in specimens S1–S5 comprise the V-notched section, the substrate, and the cladding. The fracture morphology of specimens S1, S2, S4, and S5 is analogous, as depicted in Figure 11a,b,d,e. The fracture surface’s substrate areas present a granular conglomerate-like texture, with no distinctly observed instances of either delamination or slippage. Additionally, the cladding layer appears as a dark grey pigment. However, contrary to this trend, specimen S3’s fracture morphology displays considerable delamination and slippage, emanating from the V-notch region and extending towards the substrate, encapsulating the fusion cladding. The fibrous nature of the substrate region of the fracture is apparent, and there is visible translational displacement of breccia particles in the metal as depicted in Figure 11c.
The binding area of specimens S1–S5 was expanded in a local manner, as clearly demonstrated in Figure 11a1–e1,a2–e2. A considerable number of cavities were observed in the cladding and bonding regions of specimens S1, S2, S4, and S5, which are attributed to the rapid cooling rate, resulting in the immobilization of the shielding gas during the cladding process and its subsequent failure to disperse [38]. This suggests that the cladding and substrate are not strongly bonded, indirectly. However, there are only a few minor flaws in the bonding zone and cladding layer of sample S3, suggesting a strong bond between the cladding and substrate. Notably, the coating’s pore size initially reduces and later increases as the powder mass flow rate increases. Figure 11a3–e3,a4–e4 demonstrate localized expansion of the substrate and cladding in samples S1 to S5. It can be seen that a large number of dimples and ductile tear ridges are distributed in the cladding layer area of the specimen [39]. Additionally, a significant number of streaks caused by the combination of tear ridges are distributed in the substrate area due to the specimen experiencing positive stress, fracturing through the grain, forming intergranular cracks, and generating local plastic deformation. After subjecting the coated specimens to impact, the fracture propagated from the V-notch region to the substrate and the coating. Prior to complete fracture, the substrate region tears instantaneously, the shear region is minimal, and a portion of the region exhibits brittle fracture characteristics. Nevertheless, when there is fracture in the cladding layer, it displays a certain bending deformation, followed by fracture, indicating higher fracture resistance characteristics. The fracture behavior of the specimen as a whole shows low fracture resistance.

4. Conclusions

The impact of powder mass flow rate on the coating properties of Inconel 718 cladding on a 27SiMn steel surface has been explored in this paper. The cladding layers’ cross-sectional microstructure, X-ray diffraction patterns, micro-hardness distribution, friction and wear properties, impact properties, and fracture morphology were evaluated at varying powder mass flow rates. Based on the analysis, the following conclusions were drawn:
(1) The fused coating’s width and height increase gradually as the powder mass flow rate rises. Meanwhile, the melt pool’s depth decreases with increasing powder mass flow rate. The heat-affected zone’s depth and the grain size of each coating exhibit a decreasing-then-increasing pattern as the powder mass flow rate increases. Conversely, the Ni3Fe phase volume fraction in the cladding layer shows the opposite trend to that of grain size. When the powder mass flow rate is at 20 g/min, the melt pool has a moderate depth and the cladding coating securely binds to the substrate. The cladding layer’s microstructure appears to be even, with fine grains, and the Ni3Fe phase’s volume fraction in the coating is the greatest;
(2) The micro-hardness, frictional wear properties, and impact properties of the coating exhibit identical trends as those of Ni3Fe. At 20 g/min, the coating achieves its highest micro-hardness and impact toughness (KCV) values of 328.2 HV0.5 and 45.4 J/cm2, respectively. The coating also exhibits the lowest average frictional coefficient and wear loss, measured at 0.52 and 0.066 g, respectively. The coating primarily undergoes abrasive wear as the wear mechanism, exhibiting superior mechanical characteristics and wear resistance;
(3) Bond zone defects in the fracture pattern provide indirect evidence of the bond strength between the cladding and substrate. The height difference of the fracture and the formation of dents and tear edges within the cladding region are indications that the cladding samples absorbed a significant amount of plastic work during impact fracture. At a powder mass flow rate of 20 g/min, the bonding strength between the coating and substrate is greater.

Author Contributions

The author contributions are as follows: C.Z. contributed to the conception of the study and structured the manuscript. L.S. contributed to experimental data analysis and writing of the manuscript. P.L. revised the paper. J.G. contributed to manufacturing Inconel 718 coatings by the laser cladding method. W.H. helped perform the analysis and provided reviews of the manuscript. H.Y. helped perform the analysis. All authors have read and agreed to the published version of the manuscript.

Funding

This work is supported by the Foundation of Key Laboratory of National Defense Science and Technology (No. JCKY61420052022), the National Natural Science Foundation of China (No. 52075544), the Shaanxi Province key research and development plan (NO. S2022-YF-YBGY-0437), the Fully mechanized mining hydraulic support remanufacturing key technology research “scientist + engineer” team (No. 2023KXJ-123), the Innovation Fund for graduate students of Shaanxi University of Technology: Mechanism and method of multi-energy field cooperative regulation of coating shape in laser cladding additive manufacturing (No. SLGYCX2311), the General Special Research Project of Shaanxi Provincial Department of Education (No. 22JK0312), and the Research Fund of Shaanxi University of Technology (No. SLG2123), Laser additive remanufacturing process and shape control method for functional coatings (No. SLGKYXM2306).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

Data will be made available on request.

Acknowledgments

The authors thank teachers Jun Zhou and Jiasheng Wang for their support and encouragement during the study. The contributions of other technicians involved in the various experiments are duly acknowledged.

Conflicts of Interest

The authors declare no conflict of interest.

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Coatings 13 01877 g001
Figure 1. Cladding diagram.
Figure 1. Cladding diagram.
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Figure 2. Sampling process, (a) Sampling position, (b) Impact sample geometry, (c) Metallographic sample, (d) Friction and wear sample.
Figure 2. Sampling process, (a) Sampling position, (b) Impact sample geometry, (c) Metallographic sample, (d) Friction and wear sample.
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Figure 3. Cross sections of cladding at different powder mass flow rates: (a) 16 g/min, (b) 18 g/min, (c) 20 g/min, (d) 22 g/min, (e) 24 g/min.
Figure 3. Cross sections of cladding at different powder mass flow rates: (a) 16 g/min, (b) 18 g/min, (c) 20 g/min, (d) 22 g/min, (e) 24 g/min.
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Figure 4. Localized morphology of the cladding (a–e) and microstructure of the top (a1–e1), transition zone (a2–e2), and bottom (a3–e3) of the fused coating at various powder mass flow rates. S1 (Mass flow rate of 16 g/min), S2 (Mass flow rate of 18 g/min), S3 (Mass flow rate of 20 g/min), S4 (Mass flow rate of 22 g/min), and S5 (Mass flow rate of 24 g/min).
Figure 4. Localized morphology of the cladding (a–e) and microstructure of the top (a1–e1), transition zone (a2–e2), and bottom (a3–e3) of the fused coating at various powder mass flow rates. S1 (Mass flow rate of 16 g/min), S2 (Mass flow rate of 18 g/min), S3 (Mass flow rate of 20 g/min), S4 (Mass flow rate of 22 g/min), and S5 (Mass flow rate of 24 g/min).
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Figure 5. XRD patterns of the substrate (S0) and samples of S1 (Mass flow rate of 16 g/min), S2 (Mass flow rate of 18 g/min), S3 (Mass flow rate of 20 g/min), S4 (Mass flow rate of 22 g/min), and S5 (Mass flow rate of 24 g/min).
Figure 5. XRD patterns of the substrate (S0) and samples of S1 (Mass flow rate of 16 g/min), S2 (Mass flow rate of 18 g/min), S3 (Mass flow rate of 20 g/min), S4 (Mass flow rate of 22 g/min), and S5 (Mass flow rate of 24 g/min).
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Figure 6. Micro-hardness distribution of coatings at different mass flow rates.
Figure 6. Micro-hardness distribution of coatings at different mass flow rates.
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Figure 7. Friction and wear properties of substrate S0 and coatings S1–S5: (a) friction coefficient curve, (b) wear amount.
Figure 7. Friction and wear properties of substrate S0 and coatings S1–S5: (a) friction coefficient curve, (b) wear amount.
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Figure 8. Wear morphology of substrate S0 and coatings S1–S5: (a) Substrate, (b) 16 g/min, (c) 18 g/min, (d) 20 g/min, (e) 22 g/min, (f) 24 g/min.
Figure 8. Wear morphology of substrate S0 and coatings S1–S5: (a) Substrate, (b) 16 g/min, (c) 18 g/min, (d) 20 g/min, (e) 22 g/min, (f) 24 g/min.
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Figure 9. Impact toughness of substrate S0 and cladding specimens S1–S5.
Figure 9. Impact toughness of substrate S0 and cladding specimens S1–S5.
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Figure 10. The fracture morphology of impact specimens S1–S5 and zones with typical patterns: (a–e) Three-dimensional morphology, (a1–e1) diagram of the typical zones’ locations: (I) shear zone, (II) shear lip zone, (III) subcritical crack propagation zone, and (IV) final fracture zone, (a2–e2) Area histogram.
Figure 10. The fracture morphology of impact specimens S1–S5 and zones with typical patterns: (a–e) Three-dimensional morphology, (a1–e1) diagram of the typical zones’ locations: (I) shear zone, (II) shear lip zone, (III) subcritical crack propagation zone, and (IV) final fracture zone, (a2–e2) Area histogram.
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Figure 11. Fracture morphology of impact specimens S1–S5: (a–e) Macro-morphology, (a1–e1) and (a2–e2) local amplification morphology, (a3–e3) Local SEM morphology of cladding layer, (a4–e4) Local SEM morphology of substrate.
Figure 11. Fracture morphology of impact specimens S1–S5: (a–e) Macro-morphology, (a1–e1) and (a2–e2) local amplification morphology, (a3–e3) Local SEM morphology of cladding layer, (a4–e4) Local SEM morphology of substrate.
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Table 1. Process parameters of laser cladding.
Table 1. Process parameters of laser cladding.
SamplesLaser Power
(W)		Mass Flow Rate
(g/min)		Scanning Speed
(mm/s)		Carrier Gas Flow Rate
(L/min)
S11800162210
S21800182210
S31800202210
S41800222210
S51800242210
Table 2. Cladding layer width and height, depth of the melt pool, and depth of the heat-affected zone.
Table 2. Cladding layer width and height, depth of the melt pool, and depth of the heat-affected zone.
SamplesWidth of Cladding Layer (μm)Height of Cladding Layer (μm)The Depth of the Melt Pool (μm)Depth of HAZ (μm)
S179221264364584
S279761344339571
S380691434322567
S481251548284575
S581941683222597
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MDPI and ACS Style

Zhang, C.; Shu, L.; Li, P.; Gong, J.; He, W.; Yu, H. Influence of Powder Mass Flow Rates on Wear Resistance and Impact Toughness of Inconel 718 Surface Coatings. Coatings 2023, 13, 1877. https://doi.org/10.3390/coatings13111877

AMA Style

Zhang C, Shu L, Li P, Gong J, He W, Yu H. Influence of Powder Mass Flow Rates on Wear Resistance and Impact Toughness of Inconel 718 Surface Coatings. Coatings. 2023; 13(11):1877. https://doi.org/10.3390/coatings13111877

Chicago/Turabian Style

Zhang, Chaoming, Linsen Shu, Peiyou Li, Jiangtao Gong, Wei He, and Helong Yu. 2023. "Influence of Powder Mass Flow Rates on Wear Resistance and Impact Toughness of Inconel 718 Surface Coatings" Coatings 13, no. 11: 1877. https://doi.org/10.3390/coatings13111877

APA Style

Zhang, C., Shu, L., Li, P., Gong, J., He, W., & Yu, H. (2023). Influence of Powder Mass Flow Rates on Wear Resistance and Impact Toughness of Inconel 718 Surface Coatings. Coatings, 13(11), 1877. https://doi.org/10.3390/coatings13111877

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