Effects of TiC on the Microstructure and Mechanical Properties of Four Fe-Based Laser Cladding Coatings

Abstract

This study focuses on Fe-based laser cladding coatings containing varying levels of four elements, and the objective is to investigate the influence of TiC addition on the microstructural evolution, microhardness, wear resistance, and corrosion resistance of these Fe-based composite coatings. Fe/TiC composite coatings were prepared by incorporating 20 wt.% TiC into four types of Fe-based coatings. The coatings were characterized using X-ray diffraction (XRD), optical microscopy for microstructural observation, microhardness testing, friction and wear tests, and electrochemical analysis. The results indicate that the phases of the coatings are primarily composed of α-Fe and Cr₇C₃. Upon the addition of TiC, the TiC and Ni₃Ti phases were observed in the coatings. The coatings mainly consist of columnar crystals, dendritic structures, equiaxed grains, and cellular structures, with petal-shaped TiC particles distributed within the coating matrix. TiC effectively enhances the microhardness and wear resistance of the coatings. The average microhardness of the coatings increased from 455.8 ± 20.8 HV_0.2 to 802.8 ± 41.6 HV_0.2 with TiC addition. Simultaneously, the wear rate of coating A2 decreased from 1.51 × 10⁻⁶ g/(N·m) to 1.02 × 10⁻⁷ g/(N·m), indicating an order of magnitude improvement in wear resistance. However, TiC destroys the denseness of the Fe coating, the current corrosion density increases by 28% on average, and the corrosion resistance decreases significantly.

1. Introduction

Laser cladding technology, as one of the novel surface modification techniques, utilizes high-energy laser beams to melt powders, enabling metallurgical bonding between the powder and substrate. Laser cladding technology offers advantages such as low heat output, minimal substrate distortion, and high coating adhesion strength [1,2,3]. Guo et al. [4] and others utilized laser cladding to prepare coatings on carbon steel to improve abrasion and corrosion resistance, significantly increasing the service life of steel components.

Fe-based alloy powders [5,6,7], commonly used in laser cladding, possess significant advantages, such as high microhardness and low cost. Cui et al. [8] employed laser cladding technology to fabricate Fe-based powder coatings with an average microhardness of approximately twice that of the substrate, demonstrating superior wear resistance compared to the base material. Due to its low cost and similarity in composition to structural steel, iron-based coatings are often prioritized as the material of choice for cladding [9]. Currently, numerous scholars have conducted extensive research to enhance the wear resistance and corrosion resistance of iron-based coatings, with the aim of extending their applicability to more demanding environments. Zeisig et al. [10] developed a high-microhardness and wear-resistant Fe85Cr4Mo8V2C1 coating reinforced with VC and Mo2C, exhibiting a microhardness of 829 HV and a wear rate of only 10⁻⁵ mm³/Nm. Li et al. [11] added varying amounts of Nb to martensitic stainless steel and prepared iron-based Nb (C, N) composite coatings through in situ generation methods.

Cameron Barr et al. [12] summarized the laser additive technology for repairing steel and established the link between conventional heat treatment parameters and the properties of additively manufactured steel and also clarified the mechanism of the influence of the heat input from laser additive manufacturing by comparing the other additive manufacturing and L-DED repair technologies horizontally. Ji et al. [13] prepared an iron-based amorphous coating by laser cladding on 316L stainless steel to study the effect of the amorphous phase on the coating properties, and the results show that the coating has good wear and corrosion resistance, but the increase in the amorphous phase reduces the corrosion resistance of amorphous composite coatings.

In order to face high load conditions, the wear resistance of Fe-based coatings needs to be improved. Existing studies have shown that wear resistance requires high microhardness of the material, and the microhardness and wear resistance of the coatings can be improved by adding hard carbide-reinforced phases [14,15]. Li et al. [16] found that the addition of WC significantly enhances the microhardness and performance of WC/Ni60 composite coatings. However, they also observed negative effects, such as cracking. As the mass fraction of WC increases, carbides in the cladding layer continuously increase, resulting in the production of coarser dendrites or even large crystals. Chen et al. [17] investigated the influence of TiC content on coating performance and observed that the wear resistance of the coating initially increases with increasing TiC content. However, when the mass fraction of TiC reaches 30%, it begins to decrease.

However, the addition of TiC tends to have a certain potential difference with the metal substrate, which will form local galvanic coupling corrosion and reduce the corrosion resistance of the material [18,19]. Zhou et al. [20] investigated the effect of Ni content in the fusion cladding material on the performance of the fusion cladding layer, and the results show that the addition of Ni will form a good passivation film and improve the corrosion resistance of the coating, and the corrosion resistance of the coating is increased as the content of Ni increases.

The above findings show that the addition of an appropriate amount of carbide [21,22] in the preparation of coatings using laser cladding technology can effectively and significantly enhance the organizational properties and wear resistance of the coatings. However, the mechanism of the effect of TiC on the wear resistance and corrosion resistance of Fe composite coatings is not comprehensively investigated at present. Therefore, it aims to explore the influence law of TiC addition on the organization evolution, microhardness, and wear resistance of Fe-based composite coatings, which provides an important reference for the design and application of Fe/TiC composite coatings.

2. Experimental

2.1. Coating Preparation

65Mn was chosen as the substrate, with dimensions of 150 mm × 150 mm × 50 mm, and the powder was chosen as Fe-based powder with different contents of four elements (particle size of 100–270 μm), the compositions of which are shown in

Table 1. Chemical composition of Fe-based powder (wt.%).

Number	C	B	Cr	Ni	Fe	Si	Mn	Mo	Nb
A1	0.09	0.50	15.00	7.40	Bal	0.80	0.20	1.00	0.30
A2	0.08	0.41	15.83	5.31	Bal	0.92	0.39	1.35	0.29
A3	0.20	1.13	15.84	1.89	Bal	1.19	0.15	1.03	-
A4	0.18	0.80	16.00	4.00	Bal	1.20	0.27	1.73	1.30

. The substrate was sanded smoothly with a hand-held sander and cleaned with ethanol. In view of the fact that the wear resistance of Fe-based alloy coatings is not excellent enough, the microhardness and wear resistance of the coatings can be improved by adding hard carbide-reinforcing phases. Powder selection of the above four elements was performed with different content of Fe-based powder. In each powder, 20 wt.% TiC (particle size of 150–200 μm) was added to each powder; the compositions are shown in

Table 2. Chemical composition of TiC powder (wt.%).

Element	Ti	Fe	Si	C
Content	Bal	0.19	0.004	19.21

, and the powders were mixed for 1 h using a powder mixer. The four coatings after adding TiC were B1–B4 in order. The composition of B1–B4 is shown in

Table 3. Chemical composition of Fe/TiC powder (wt.%).

Number	C	B	Cr	Ni	Fe	Si	Mn	Mo	Nb	Ti
B1	3.91	0.40	12.00	5.92	Bal	0.64	0.16	0.80	0.24	16.12
B2	3.90	0.33	12.66	4.25	Bal	0.74	0.31	1.08	0.23	16.12
B3	4.02	0.90	12.67	1.51	Bal	0.95	0.12	0.82	—	16.12
B4	4.00	0.64	12.80	3.20	Bal	0.96	0.22	1.38	1.04	16.12

.

Using an IPG-6000W fiber laser based on the existing literature research and prior work, the coating was prepared under the process parameters of a laser power setting of 2000 W, a scanning speed of 10 mm/s, a powder feeding rate of 1.6 r/min, a lap rate of 60%, a spot diameter of 2.5 mm, and a protective gas argon flow rate of 15 L/min. Figure 1 shows the laser melting Fe-based coatings of A1–A4, and Figure 2 shows the laser melting Fe/TiC laser melting composite coatings of B1–B4. In Figure 1, it can be observed that none of the four coatings exhibit macroscopic cracks. This is attributed to the good ductility of the Fe-based substrate and the moderate content of Si and B, which facilitate deoxidation and self-fluxing. These elements possess good slag-forming abilities, ensuring the formation of a well-shaped molten pool during the laser cladding process, thereby resulting in well-formed coatings.

2.2. Microstructural Characterization and Wear Testing

Physical phase changes of four Fe-based coatings and coatings with TiC addition were observed using XRD (Pulsar XD-3, Beijing, China), and 10 × 10 mm² specimens and friction specimens with a diameter of Φ25 mm² were cut using wire cutting. Using an Vickers microhardness tester (HV1000z, Shanghai, China) with a load of 200 g and a holding time of 10 s, the test was carried out at intervals of 0.25 mm along the cross-section from the coating to the substrate, and the distance from each point to the coating surface was recorded. In order to minimize the interference of residual TiC microhardness relative to microhardness, the microhardness test was performed to avoid the visible hard reinforcement phase. The microstructure of the Fe-based coatings after 4% nitric acid alcohol etching and the microstructure of the Fe/TiC composite coatings after aqua regia etching were observed using a 4XC model after sanding and polishing with sandpaper.

Using a friction tester (CHMT23, Beijing, China), the coatings were rubbed with a silicon nitride ball at 500 rpm with a load of 50 N for coatings for 20 min, the difference in mass before and after abrasion was measured using an electronic balance (FA2004B, Shanghai, China) with an accuracy of 0.1 mg, and the average value was taken by repeating the experiment three times. This study utilizes a Carl Zeiss scanning electron microscope (EVO18, Oberkochen, Germany) to examine the wear morphology. The corrosion resistance of the coating was tested using an electrochemical workstation (Ch660E, Shanghai, China) with a platinum electrode and a reference electrode of calomel using 3% sodium chloride.

3. Results and Discussion

3.1. Physical Phase Analysis

Figure 3 and Figure 4 show the physical phase diagrams of the coatings determined by comparing the diffraction peaks according to the XRD energy spectrum database [23]. The main component of the iron-based alloy melting and coating powder is Fe, followed by higher contents of Cr and Ni, a certain amount of Mo, Mn, and Si, and other elements (see Table 1). The laser heat input source melts the alloy powder to form a molten pool in which the above alloying elements undergo chemical metallurgical reactions under the action of the laser. As can be seen in Figure 3, the composition of the Fe-based coatings with four different elements is mainly composed of α-Fe, which corresponds to the elemental content in Table 1. The other phases are also different from the elemental content. As can be seen in Figure 3, the X-ray diffraction spectra of A1 and A3 are similar, generating the hard phase Cr₇C₃, while A2 and A4 are similar, generating Cr_0.03Fe_0.97. The main phases of the four coatings are composed of α-Fe and residual TiC, generating Ni₃Ti, as well as the hard phase Cr₇C₃ generated by Cr and C. The main phases of the four coatings are composed of α-Fe, residual TiC, generating Ni₃Ti, as well as Cr and C, generating Cr₇C₃.

In Figure 4, it can be seen that the diffraction peaks of TiC for B2 and B3 are the strongest, which indicates that B2 and B3 have the highest TiC content on the coating surface, whereas the weaker diffraction peaks of TiC for B1 and B4 indicate that their surfaces have less TiC content. This is due to the lower density of TiC particles (4.9 g/cm³) compared to the Fe-rich substrate (7.8 g/cm³) in the molten pool generated by the melting of the powder by the laser beam, followed by the decomposition of the TiC particles while slowly floating up to the top of the substrate in the molten pool, while the lower cooling rate leads to the increase in the localized temperature in the bonding area of the melting zone with the substrate. As a result, TiC particles are more likely to gather at the top of the fused cladding layer. The residual TiC content is higher in B2 and B3. As can be seen from Table 1 powder elements, B4 has a higher content of C and Ni, and enough Ni and Ti in B4 combine to generate the hard phase Ni₃Ti, which promotes the decomposition of TiC, and it is also possible that the uneven mixing of TiC particles and Fe-based alloy powder particles leads to this phenomenon.

3.2. Microstructure Analysis

Figure 5 shows the microstructure of the middle layer of the A1–A4 fused cladding; the fused cladding is overall flat, and no defects such as pore cracks are observed. Normally, the molten pool cools faster due to the high temperature and the low temperature of the substrate. The grains do not have enough time for adequate nucleation and growth. Since the growth direction of the crystals tends to grow along the grain direction with the lowest energy, the bottom of the molten pool is mostly columnar crystals [24]. In Figure 5, it can be seen that the A1 coating is mainly converted from dense cellular crystals to columnar crystals, and the growth direction of the branch crystals is at a certain angle to the interface. It can be observed in Figure 5 that the A2–A4 coating is mainly composed of columnar crystals and dendrites, and it can be observed in Figure 5A2 that the columnar crystals are gradually converted into dendrites along the solidification direction (from right to left).

According to solidification theory, the stability factor (G/R) is an important parameter between the solid–liquid interface, determining the morphology of the solidification structure. Here, G represents the magnitude of the temperature gradient, while R indicates the speed of solidification [25]. During the laser cladding process, as the laser continues to heat up, the metal powder absorbs heat and transforms into a molten state. At this time, the cooling rate of the upper layer of the melt pool is slow, while the substrate at the bottom of the melt pool dissipates heat quite rapidly. In addition, the uppermost layer of the molten cladding layer dissipates most of its temperature in the direction of the substrate as the air dissipates heat, resulting in a relatively high-temperature gradient G. The solidification rate R tends to decrease from the substrate to the upper layer of the molten cladding layer, and the microstructure of the upper layer of the molten cladding layer changes to finer cellular crystals with the change in the G/R value. From the region where the cooling rate of the melt pool is accelerated, the solidification rate is increasing, contributing to the formation of fine cellular crystals.

From the comprehensive analysis in Figure 3, it is observed that enhanced phases are generated in A1–A4, as evident from the XRD results. However, complete morphology cannot be observed solely through microscopic observation. The enhanced phase in the A1 coating primarily exhibits an irregular block-like distribution, while in A2–A4, it is predominantly distributed in a dispersed strengthening manner. Some black holes can be observed in Figure 5, which are remnants of fine strengthening phase particles precipitated within the coating and corroded. This phenomenon arises due to the presence of trace amounts of carbon (C) in the metal powder, leading to the precipitation of a certain quantity of fine strengthening phase particles within the coating. These particles detach during the metallographic corrosion process. In comparison with other coatings, the size of carbide particles within A4 is further reduced, exhibiting a more uniformly dispersed distribution. This phenomenon is also observed in the research conducted by the Ye Jie Liang team.

Figure 6 shows the microstructure (B1–B4) and the elemental distribution of the coating cross-section (B5–B8) of the four TiC-incorporated fused cladding layers. The fused cladding layers were overall flat and no defects, such as pore cracks, were observed. From the elemental distribution diagram, it can be seen that the irregular lumps are rich in Ti and C, and there is almost no Fe, which can be judged as a TiC hard phase. In Figure 6, a large number of coarse petal-like TiC particles can be observed, as well as fine particles of TiC. Among them, the TiC hard phases in B2 are the most densely populated, while those in B3 are slightly weaker, consistent with the phase diagram of the coatings in Figure 4. The shapes of the enriched large TiC particles vary, and it is observed that larger TiC hard phase particles contain central pits. This is attributed to the simultaneous decomposition and precipitation of TiC hard phases into smaller fragments. Therefore, fine black TiC dots can still be observed in B1 in Figure 6, indicating a more uniform distribution. The second phase of small TiC particles is relatively high in B1 and B2 coatings, exhibiting a cellular structure around the particles. Additionally, with the partial dissolution of TiC, the concentrations of Ti and C increase, enhancing the stability of the coatings after solidification.

B1 and B4 are rich in Ni and Cr. The decomposition of TiC leads to the formation of dispersed hard phases, where the decomposed Ti combines with Ni to produce the hard phase Ni₃Ti, while Cr combines with C to generate hard phases, such as Cr₇C₃. When the content of decomposed TiC is high, it can precipitate again, forming fine TiC particles. In contrast, B3 and B2 have lower Ni content and higher residual TiC content.

3.3. Microhardness Analysis

Figure 7 shows the microhardness distribution of four Fe-based coatings. The microhardness trends of these four Fe-based coatings in the clad layer are similar. In the transition from the clad layer to the heat-affected zone (HAZ), the microhardness fluctuates around 400 HV_0.2 and then slightly decreases to around 150 HV_0.2 in the HAZ, while the microhardness of the 65Mn substrate is approximately 150 HV_0.2. The entire coating can be divided into the clad zone (CZ), the HAZ, and the substrate (Sub). The average microhardness of the fused cladding layer of A1 was 420.8 ± 18.8 HV_0.2, the average microhardness of the fused cladding layer of A2 was 424.8 ± 37.4 HV_0.2, the average microhardness of the fused cladding layer of A3 was 442.7 ± 21.8 HV_0.2, and the average microhardness of the fused cladding layer of A4 was 455.8 ± 20.8 HV_0.2. Observing the curves, it can be noted that there is a relatively small difference in microhardness among the four powders, which is also reflected in the microstructures shown in Figure 5.

This is because when the laser power is appropriate, the heat input is low, the tail of the molten pool is short, and the solidification rate is fast; the structure of the clad layer does not grow, and the reinforcing phases generated are dispersed as fine particles in the matrix of the clad layer, resulting in a higher quality clad layer (see Figure 1). In Figure 7, it can be observed that the microhardness in the middle region of the clad layer is higher than that in the upper region and significantly higher than the surface microhardness. Although the grains on the surface of the clad layer are relatively small, the microhardness is still lower. The cellular crystals in the middle of the clad layer hinder the movement of dislocations, which is the main reason for the increase in microhardness. Similar conclusions were also drawn by Shi et al. [26]. The microhardness order of the coatings is A4 > A3 > A2 > A1.

The residual TiC particles in the coating are black in color, and their microhardness reaches more than 1000 HV, which will interfere with the microhardness test results and subsequent analysis, so avoid TiC particles when testing microhardness. Figure 8 illustrates the microhardness distribution of four coatings. In Figure 8, it can be observed that the microhardness distribution in the clad layers of these four laser cladding coatings is relatively uniform. This indicates that after removing the residual TiC particles, the microstructure of the coatings is uniform, which allows for better encapsulation of TiC particles. It is this cement-like structure that further enhances the overall performance of the coatings, enabling them to firmly grasp the hard phases under high-load conditions, thereby reducing the likelihood of coating delamination and exhibiting improved microhardness and wear resistance.

The average microhardness of the B1 coating was 555.3 ± 20.1 HV_0.2, and the average microhardness of the B2 coating was 467.3 ± 20.1 HV_0.2. The average microhardness of the B3 coating was 729.8 ± 48.6 HV_0.2, and the average microhardness of the B4 coating was 802.8 ± 41.6 HV_0.2, respectively, with a gradual increase in microhardness from B1 to B4. The microhardness increases significantly towards the coating–substrate interface, with coatings B3 and B4 having similar microhardness values exceeding 700 HV_0.2, while the performance of B2 is poor, with B4 exhibiting the highest microhardness. Combining the phase composition in Figure 4 and the microstructure of the coatings in Figure 6, it can be inferred that B2 has a higher content of residual TiC, with fewer hard phases, such as Cr₇C₃ and Ni₃Ti, in the coating. This deliberate avoidance of TiC particles during microhardness testing is one of the reasons for the poor microhardness performance of B2. Several significant fluctuations in microhardness values are still present in the curves. For example, the microhardness of the B3 coating undergoes a significant abrupt change between distances of 1.0 to 1.3 mm from the coating surface. This may be attributed to the presence of TiC particles hidden beneath the detection area and an increase in the G/R ratio, resulting in a slower solidification rate compared to the upper layers, which is caused by the overly concentrated distribution of dispersed hard phases. Overall, the microhardness range of the coatings, from highest to lowest, is B4 > B3 > B1 > B2, which corresponds to the trends observed in the phase composition and microstructure of the coatings.

3.4. Friction and Wear Behavior Analysis

Table 4. Friction behavior of the coating.

Number	Average Coefficient of Friction	Mass Wear Rate g/(N·m)
A1	0.57	2.19 × 10−6
A2	0.50	1.94 × 10−6
A3	0.57	1.84 × 10−6
A4	0.55	1.51 × 10−6
B1	0.54	0.97 × 10−7
B2	0.53	1.02 × 10−7
B3	0.52	0.83 × 10−7
B4	0.43	0.82 × 10−7

presents the average friction coefficients of the coatings, sorted in ascending order: A4 < A3 < A1 < A2. Figure 9 illustrates the fluctuation of friction coefficients over time for the four types of deposited layers. The wear process is generally divided into initial wear and stable wear stages. Within 0–300 s, the friction coefficient between the substrate and the coating sharply increases and then slightly decreases. This phenomenon occurs because during the initial contact between the sphere and the coated surface, the contact area is small, and the stress is high. As wear progresses over time, the contact area is compressed and sheared, and a large number of wear particles are generated, causing “plowing” on the coating surface. Simultaneously, the worn area on the sphere increases, leading to a rapid increase in the friction coefficient. With time, as the surface becomes smoother due to wear, the system enters a stable wear stage, where the friction curve becomes relatively stable. However, after the mid-term stable stage, significant fluctuations in the friction curve reappear. This is attributed to the complete wear of the hard phase within the coating by the small sphere, resulting in an increase in the surface roughness of the substrate, which, in turn, affects the fluctuation of the friction coefficient. The friction coefficient is the ratio of friction force to normal load and is mainly determined by factors such as surface roughness, load, sliding velocity, temperature, and friction pair materials.

To better assess the wear resistance of the coatings, we introduce the mass wear rate as a criterion. The mass wear rate is calculated as the ratio of the difference in the mass of the coating before and after friction wear testing to the product of the load and sliding distance, i.e., wear mass difference/friction work. This metric provides a more intuitive reflection of the wear resistance of the coatings [27].

The mass wear rate is as follows:

Mass Wear Rate: A = Δm/(S·P)

(1)

where A is the wear rate g/(N·m), Δm is the mass difference g before and after friction, S is the friction distance m, and P is the added load N.

According to tribological theory, when the general microhardness is below 800 HV, the wear resistance of the coating is positively correlated with the microhardness, and when the microhardness of the coating reaches more than 800 HV, the coating is prone to plastic deformation; the anti-shear ability becomes poorer, and the microhardness of the coating is negatively correlated with the wear resistance.

In Figure 9, it can be observed that in the first 200 s, the fluctuation trends of the four coatings are similar, with the friction curves rapidly rising and then stabilizing. This is because during the laser cladding process, the Fe-based coatings are rich in elements, such as Cr and Ni, which combine with Ti and C to generate hard phases, such as Ni₃Ti and Cr₇C₃, and these hard phases are deposited in the molten pool. When the Si₃N₄ balls undergo abrasion in the 0–200 s period, the wear sharply increases due to initial contact with the hard cladding. After 600 s, the wear of the upper hard phases reaches equilibrium, and the friction coefficient gradually stabilizes. Similarly, in Figure 9, in the first 200 s, the fluctuation trends of the four coatings are similar, with the friction curves rapidly rising and then stabilizing. This is because during the laser cladding process, the Fe/TiC composite coatings are rich in elements, such as Cr and Ni, which combine with Ti and C to generate hard phases, such as Ni₃Ti and Cr₇C₃, and these hard phases are deposited in the molten pool. When the Si₃N₄ balls undergo abrasion in the 0–200 s period, the wear sharply increases due to initial contact with the hard cladding. After 600 s, the wear of the upper hard phases reaches equilibrium, and the friction coefficient gradually stabilizes.

Based on the mass wear rate in Figure 10, the mass wear rates of coatings A1 to A4 are all on the order of 10⁻⁶ g/(N·m), among which the A4 coating is the lowest at only 1.51 × 10⁻⁶ g/(N·m). Solely based on the mass wear rate, the wear resistance of the four coatings increases sequentially from A1 to A4. The friction coefficients of A1 to A4 are close, and a phenomenon similar to the composition of the four powders in Figure 3. In Figure 10, the wear rates of the four coatings, A1 to A4, also show a decreasing trend, indicating that the wear resistance of the A4 coating is optimal. This is supported by the microstructure images in Figure 5. The friction coefficient of A4 is around 0.55, with a wear rate of 1.51 × 10⁻⁶ g/(N·m), while the friction coefficient of A3 is around 0.57, with a wear rate of 1.84 × 10⁻⁶ g/(N·m). Overall, the wear resistance of A3 is also close to that of A4. The ranking of wear resistance of the coatings is A4 > A3 > A2 > A1. The performance of wear resistance is consistent with the microhardness of the coatings.

As can be seen in Table 4, the average friction coefficients of the B1–B4 coatings are 0.54, 0.53, 0.52, and 0.43. The average friction coefficients of the B1, B2, and B3 coatings are close to each other, while the average friction coefficient of the B4 coating is only 0.43. From the perspective of the mass abrasion rate in Figure 10, the mass abrasion rates of the B1–B4 coatings are very close to each other, and although there are fluctuations, the differences between the four coatings’ wear resistance are not very large; from the mass abrasion rate alone, the B4 coating has the lowest mass abrasion rate of only 0.82 × 10⁻⁷ g/(N·m). Within the range of the four coatings, the difference in wear resistance is not large, and from the mass wear rate alone, the B4 coating has the lowest mass wear rate of only 0.82 × 10⁻⁷ g/(N·m), and the B4 coating has the smallest wear fluctuation. The microhardness has a high positive correlation with the abrasion resistance of the laser melting cladding, and the addition of TiC can significantly improve the abrasion resistance of the coating. Since the microhardness of TiC is higher than that of Si₃N₄ spheres, the scratches in the area where TiC is present in the wear test are shallower. TiC is tightly bonded to the coating and does not easily come off. The TiC in the coating slows down the formation and extension of scratches.

Figure 11 illustrates the wear morphology of four Fe-based coatings. It can be clearly observed in Figure 11 that the worn surface of the A1 coating exhibits areas of spalling, with the direction of spalling roughly parallel to the direction of the ball sliding, and the worn surface presents concave structures. The scratch depth on the sample surface is more pronounced, as are the spall pits. It can be seen that plastic deformation occurred in the A1 sample after wear. The results indicate that its main wear mechanisms are adhesive wear and abrasive wear [28].

On the worn surface of A2, small abrasives and shallow, closely spaced microplows parallel to the direction of motion are observed. No obvious adhesion or spalling phenomena are observed. The bottom surface remains relatively flat, with only slight wear. The predominant wear mechanism is abrasive wear. It can be seen that the wear resistance of the A2 coating is superior to that of the A1 coating. After wear, the A3 coating exhibits smaller spall pits, indicating lighter adhesive wear and abrasive wear. This is because the microstructure of A3 is finer, and this superior wear resistance compared to the former is can also be demonstrated from the phase diagram in Figure 3 and the microstructure diagram in Figure 5. The worn surface of the A4 coating shows large and slight plow marks but few spall pits, with the main wear mechanism being adhesive wear. This is due to the more uniform structure of A4, which is less prone to hard phase spalling and resultant abrasive wear. This is also confirmed in the microstructure diagram in Figure 5.

Figure 12 shows the wear morphology of four coatings with 20 wt.% TiC addition after 20 min of friction under a 50 N load. Compared with coatings containing a large number of hard phases, there is only a small amount of hard phase (such as Ni₃Ti) in the coatings, and a considerable number of soft phases is present, leading to a relatively reduced wear resistance. In the wear test, the test block is pressed against the disc under a 50 N load, and the disc operates under conditions of 500 r/min. Sliding dry friction occurs between the test block and the ball. During the wear process, scratches form on the surface of the test block. As the wear time increases, the number and depth of scratches increase. The TiC in the coating can slow down the formation of scratches, reduce the severity of wear, and minimize the occurrence of spalling and the dropping of wear debris caused by scratches. With the increase in hard phase content in the coating, the wear mass loss of the specimen decreases at the same wear time. This is also due to the reduction in spalling caused by wear. Conversely, if the soft phase in the coating gradually increases, the opposite result will be obtained. The main wear mechanisms of the coating are adhesive wear and abrasive wear. In Figure 12, the number of spall pits in the B3 coating is significantly reduced compared to B1 and B2, and the grooves are relatively shallow, indicating a decrease in adhesive wear of the coating. This suggests that the main wear mode has shifted to abrasive wear. From the wear marks of B4 in Figure 12, it can be seen that the wear resistance of B4 powder with TiC addition is further increased, and the grooves on the wear surface are significantly deepened. It can be observed that the main wear mechanism is abrasive wear. Overall, in terms of friction behavior, the wear resistance of the coatings ranks as follows: B4 > B3 > B1 > B2. This result is consistent with the above results of phase composition, microhardness, and TiC content.

3.5. Corrosion Behavior Analysis

The polarization curve test results are shown in Figure 13, and the corresponding statistical data for the average corrosion potential (E_corr) and corrosion current density (I_corr) are shown in

Table 5. Coating corrosion potential and corrosion current density.

Number	Corrosion Potential (V)	Corrosion Potential Density (A·cm−2)
A1	−0.570	3.28 × 10−5
A2	−0.515	5.18 × 10−5
A3	−0.571	2.59 × 10−5
A4	−0.553	3.71 × 10−5
B1	−0.575	4.93 × 10−5
B2	−0.476	3.55 × 10−5
B3	−0.518	4.01 × 10−5
B4	−0.456	4.37 × 10−5

. The corrosion potential represents the material’s tendency to corrode. The higher the potential and the lower the corrosion current density, the less likely corrosion will occur. From the comprehensive results in Figure 13 and Table 5, it can be seen that the corrosion potential of A3 is −0.571 V, and the relative lowest corrosion current density is 1.09 × 10⁻⁵ A·cm⁻². However, it is not accurate to judge the difference in corrosion resistance based solely on the polarization curve and corrosion current density. In order to better assess the corrosion behavior of the coatings, electrochemical impedance spectroscopy tests were conducted, and the results are shown in Figure 14. It can be clearly seen in Figure 14 that the capacitive arc radius of A3 is the largest, indicating that the charge transfer impedance on the surface of the A3 coating is greater. This means that the diffusion of ions in the solution is more difficult, and, therefore, the corrosion resistance of A3 is better. The ranking of corrosion resistance of the coatings is as follows: A3 > A4 > A2 > A1. This corresponds to the results in Table 5 and Figure 5.

In alloy systems, under the conditions of laser cladding with high cooling rates, Si and Mn segregate at grain boundaries, reducing the temperature of the eutectic austenite and refining the grains, thereby improving corrosion resistance. Additionally, the coatings contain corrosion-resistant alloying elements, such as Cr and Ni, to further enhance corrosion resistance. Overall, A3 has the highest relative content of Cr and Ni, resulting in the best corrosion resistance. The trends in corrosion resistance performance correspond to the contents of Cr, Ni, and Mn.

The corrosion performance of each Fe/TiC cladding layer was tested through electrochemical experiments, and the corresponding statistical average values of corrosion potential (E_corr) and corrosion current density (I_corr) are presented in Table 5 and Figure 13. Combining the results in Table 5 and Figure 14, it can be observed that after the addition of TiC, the corrosion potential of B4 is relatively highest at −0.456 V, while the lowest corrosion current density is exhibited by the B2 coating at 3.55 × 10⁻⁵ A·cm⁻². The impedance spectroscopy plot in Figure 14 distinctly illustrates the ranking of corrosion resistance performance among the coatings: B4 > B2 > B1 > B3.

Comparing the coatings after the addition of TiC to the previous Fe-based coatings, all coatings exhibited a decrease in performance. For instance, the corrosion potential of the B1 coating increased from −0.570 V to −0.575 V, while the corrosion current density increased from 3.28 × 10⁻⁵ A·cm⁻² to 4.93 × 10⁻⁵ A·cm⁻². This decrease in performance is attributed to the disruption of the stable passive film formed on the surface of Fe-based coatings by TiC, despite the inherent good chemical erosion resistance of the TiC ceramic phase. However, the B2 coating showed improvement in both corrosion potential and corrosion current density after the addition of TiC. This improvement can be attributed to the originally poor corrosion resistance of the A2 coating without TiC, while the addition of TiC particles with better corrosion resistance dispersed in the coating enhanced the corrosion resistance of the B2 coating.

4. Conclusions

This paper prepared four Fe-based laser cladding coatings (A1–A4) and characterized them using X-ray diffraction phase analysis, microscopic morphology analysis, wear resistance testing, post-wear surface morphology analysis, and corrosion resistance testing. The following conclusions were drawn:

(1)

There were no obvious defects on the surfaces of the four Fe-based coatings, and the coating structures were uniformly dense. The main phases of all four coatings were dominated by Fe. The X-ray diffraction spectra of A1 and A3 were similar, generating the hard phase Cr_0.03Fe_0.97, while A2 and A4 were similar, producing the hard phase Cr₇C₃.

(2)

There was a positive correlation between coating microhardness and wear resistance. The average microhardness of the four coatings exceeded three times that of the substrate, with A4 having the highest fusion layer microhardness at an average of 455.8 HV_0.2. The coefficient of friction for A4 was 0.55, with a wear rate of 1.51 × 10⁻⁶g/(N·m), indicating that this coating exhibited the best wear resistance among the four Fe-based coatings.

(3)

The content of Cr and Ni was positively correlated with the corrosion resistance of the coatings. The coating with the highest content of Cr and Ni, A3, had the lowest corrosion current density of 1.09 × 10⁻⁵A·cm⁻², indicating its optimal corrosion resistance.

(4)

Excess Cr and C played a role in the solid solution strengthening of the coatings, with Cr also contributing to grain refinement. Excessive Ni content led to a decrease in coating microhardness and wear resistance, while an increase in Mo content improved the mechanical properties of the coatings. The content of Cr and Mo in different coatings was highly correlated with their microhardness and wear resistance.

(5)

After the addition of TiC, the average microhardness of the coating was increased from 455.8 ± 20.8HV_0.2 to 802.8 ± 41.6HV_0.2, while the A2 wear rate of the coating was reduced from 1.51 × 10⁻⁶g/(N·m) to 1.02 × 10⁻⁷g/(N·m), and the abrasion resistance was improved by one order of magnitude.

(6)

After adding TiC, the stable passivation film generated on the surface of the Fe-based coating was destroyed by TiC, the corrosion resistance was reduced, and the current corrosion density increased by 28% on average. The corrosion resistance of the coating after adding TiC was ranked B4 > B2 > B1 > B3.