Microstructure and wear resistance of in-situ TiC-reinforced low chromium iron-based hardfacing alloys

Microstructure and wear resistance of in-situ TiC-reinforced low chromium iron-based hardfacing alloys | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Research Article Microstructure and wear resistance of in-situ TiC-reinforced low chromium iron-based hardfacing alloys Zhixiang Tong, Wei Shao, Chengxing He, Dingyong He This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-3530208/v1 This work is licensed under a CC BY 4.0 License Status: Under Review Version 1 posted 5 You are reading this latest preprint version Abstract Strengthening low chromium iron-based alloys with multiple alloy elements is a crucial strategy for developing "casing-friendly" hardbanding materials. In this paper, four kinds of low chromium iron-based hardfacing alloys with different Ti contents were prepared by Flux-Cored Arc Welding technology and the effect of Ti content on the microstructure, hardness and wear performance under the abrasive wear and dry sliding wear were analyzed. The in-situ TiC-reinforced iron-based hardfacing alloys exhibited microstructure comprising austenite, martensite, eutectic carbides M 7 C 3 and in-situ TiC particles. The in-situ formation of TiC particles consumed the carbon in the alloy and thus raised the martensite start temperature, resulting in a reduction in the mass fraction of austenite and eutectic carbides M 7 C 3 and an increase in the martensite. Among four hardfacing alloys, the alloy with 5 wt. % Ti exhibited the highest hardness (836.3 HV0.2), the lowest weight loss and the lowest wear rate due to its highest martensite mass fraction (69.97 wt. %). The matrix with in-situ TiC particles and martensite effectively resisted the cutting of abrasive grains and the wear mechanism developed from the microploughing in 0Ti alloy to microcutting in in-situ TiC-reinforced iron-based hardfacing alloys in abrasive wear tests. In dry sliding wear tests, TiC particles were released and slid on the surface of the alloys, resulting in the formation of grooves, while the formation of the tribochemical reaction layers contributed to a reduction in the friction coefficient and wear rate. This study provides a theoretical foundation for the development of "casing-friendly" hardbanding materials. Flux-Cored Arc Welding In-situ TiC Wear resistance Low chromium iron-based hardfacing alloys Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Figure 8 Figure 9 Figure 10 Figure 11 Figure 12 Introduction Hardbanding is essential for protecting drill pipe joints against abrasion from casing and rock particles. As the environments of oil and gas extraction become increasingly severe, the requirements for hardbanding materials also escalate. [ 1 , 2 ]. The development of hardbanding materials has transformed from high hardness and wear resistance to "casing-friendly" materials that possess good wear resistance, low friction coefficient, and anti-spalling properties[ 3 , 4 ]. One important deposition technique for hardbanding is Flux-Cored Arc Welding (FCAW). Extensive use of FCAW has been found in industries such as aerospace, coal, oilfield, and agriculture due to its high deposition rate, high productivity, simplicity of equipment, and field processability[ 5 – 7 ]. Therefore, the development of "casing-friendly" flux-cored wires for FCAW offers an economical solution to address the wear and cost challenges in hardbanding. High chromium white cast irons (12–30 wt. % Cr, 2.0–3.6 wt. % C) are an economic and traditional hardbanding material known for high hardness and excellent wear resistance[ 5 , 8 – 11 ]. However, hypereutectic white cast irons with coarse primary carbides M 7 C 3 (M is generally Cr, Fe and a few other alloy elements) are susceptible to cracking under impact conditions[ 7 ], and hypoeutectic white cast irons exhibit brittle fracture of carbides under high-stress wear conditions[ 11 ]. These cracking phenomena greatly limit the service life of hardbanding. Additionally, these materials have a high friction coefficient, which can cause significant casing wear. On the other hand, low chromium iron-based alloys (Cr < 12 wt. %, C < 2.0 wt. %) are known for excellent fracture strength and resistance to thermal fatigue cracking[ 12 – 14 ]. However, low hardness and low wear resistance limit their suitability for hardbanding applications. To meet the requirements for hardness, wear resistance, and anti-cracking, a new research direction for "casing-friendly" hardbanding involves the use of multi-alloy elements to reinforce low chromium iron-based alloys by introducing Ti, Nb, B and other alloy elements into original powder design. As a result, carbides such as TiC, NbC, and B 4 C are precipitated into the weld pool[ 6 , 15 ]. These carbides can act as heterogeneous nucleation cores, effectively refining the microstructure. The hardfacing alloys with refined microstructure possess high hardness and strength. Simultaneously, the carbides dispersed on the surface of the hardfacing alloy can significantly improve wear resistance. Among the carbides, TiC is commonly used as the reinforcing phase in metal matrix composites, including iron-based, nickel-based, aluminum-based, cobalt-based, and titanium-based composites[ 14 , 16 – 18 ]. Shi-Li Shu et al.[ 14 ] incorporated TiC nanoparticles into high-Cr hot work die steel, and they found nano-TiC provided nucleation sites for γ-Fe dendrites and prevented γ-Fe growth during solidification. As a result, the steel with TiC nanoparticles exhibited higher strength and toughness. Yan Liu et al.[ 18 ] prepared aluminum-based composite coatings reinforced with TiC particles using direct laser deposition. They found that the uniformly distributed TiC particles effectively reduced porosity defects and stopped the abrasive grains from pressing into the coating, thus enhancing the wear resistance of the coating. A. Bedolla-Jacuinde et al.[ 19 ] found that Ti could be used to increase the hardness without affecting fracture toughness by decreasing the volume fraction of eutectic carbides in the 16 wt. % Cr, 2.5 wt. % C white cast iron. However, to our best knowledge, there is limited research on the influence of Ti content on the microstructure and mechanical properties of low chromium iron-based hardfacing alloys. In this paper, four kinds of low chromium iron-based hardfacing alloys with different Ti contents were prepared by FCAW technology, and the effect of Ti content on the microstructure, hardness, and wear performance under the abrasive wear and dry sliding wear was analyzed. Experimental process 2.1. Material preparation Metal-cored flux-cored wires were developed by encasing flux-cored powder with cold-rolled annealed carbon steel strip. The strip, with 0.5 mm thickness and 12 mm width, was shaped into a U-shape, infused with 200–400 mesh flux-cored powder, and then compressed to a circle with a diameter of 1.6 mm. The powder filling coefficient ranged from 24–28%. The composition of the flux-cored powder is outlined in Table 1 . Different Ti contents (0, 3, 5 and 7 wt. %) were regulated by the addition of different ferrotitanium contents (0, 16, 26 and 34 wt. %) respectively, as shown in Table 1 . The obtained alloys are referred to as 0Ti, 3%Ti, 5%Ti and 7%Ti alloys in the paper. Table 1 Chemical composition (wt. %) of the metal-cored flux-cored wires Raw Powder Content (wt. %) Ferrotitanium (70 wt. % Ti) 0, 16, 26, 34. High carbon ferrochrome (68 wt. % Cr, 8 wt. % C) 40–50 Ferromolybdenum (60 wt. % Mo) 5–10 Ferrovanadium (50 wt. % V) 2–6 Electrolytic manganese (pure) 6–10 Graphite (pure) 2–6 Ferrosilicon (75 wt. % Si) 8–12 Aluminum (pure) 0–4 Iron (pure) Bal. The hardfacing process was conducted on a Q235 steel plate measuring 150 mm × 100 mm × 100 mm, utilizing a Fronius TPS5000 digital multifunctional welder operated by an ABB-IR60 six-axis robot, as shown in Fig. 1 (a). The chemical composition of the substrate is outlined in Table 2 . Before hardfacing, the substrate underwent mechanical polishing and cleaning in ethanol. Figure 1 (b) illustrates the schematic diagram of the hardfacing process, with the interlayer temperature maintained between 200°C and 300°C. To minimize dilution effects, three layers were deposited and the corresponding hardfacing parameters are provided in Table 3 . Table 2 Chemical composition (wt. %) of Q235 steel plate C Si Mn P S Fe 0.074 0.031 0.101 0.011 0.013 99.710 Table 3 The welding parameters Parameter Value Arc voltage 22–24 V Welding current 280–300 A Electrode polarity positive Wire feed speed 8.0 m/min Weave width 20 mm Welding speed 3 mm/s Stick-out 15–20 mm Shield gas 80% Ar + 20% CO 2 2.2. Material characterization Metallographic specimen measuring 10 mm × 10 mm × 10 mm were obtained from the hardfacing layer. These specimen were ground, polished, and etched by a solution that contains 3 mL of 68% nitric acid, 15 mL of 38% hydrochloric acid, 3 g of ferric chloride, and 50 mL of deionized water. The phase structure was examined via X-ray diffraction (XRD, BRUKER, D8 ADVANCE, German) with Cu-Kα radiation, with scanning angles ranging from 20 to 95° at room temperature. The microstructure and phase compositions were analyzed by a scanning electron microscope (SEM, FEI, QUANTA-650, Netherlands) equipped with an energy dispersive spectroscopy (EDS, OXFORD, X-MAX, Britain). Additionally, optical emission spectrometry (OES, BRUKER, Q4, German) was employed for chemical composition analysis. The microhardness was measured by a digital microhardness tester (HXD-1000A, China) with 1.96 N load applied for 10 s. 2.3. Wear tests The abrasive wear tests were conducted at room temperature by a wet sand rubber wheel tester machine (MLS-225, China) and the test method was conducted according to ASTM G105, as depicted in Fig. 1 (c). The wear specimen were cut from the top layer of the hardfacing alloys with a dimension of 57 mm × 25mm × 10 mm. The rubber wheel had a diameter of 176 mm and a hardness of 60 Shore A, rotating at a speed of 240 rpm. The abrasive slurry consisted of 1 kg water mixed with 1.5 kg quartz sand with the size in the range of 210–380 µm. The wear test lasted 30 min and a load of 100 N was applied. Before the test, a pre-grinding process was performed for 5 min to eliminate the impact of surface roughness. The specimen were cleaned in ethanol by an ultrasonic cleaner for 3 min before and after the tests, and their weight was measured to an accuracy of 1 mg on an electronic balance (Sartorius, BS224S, China). The worn surface morphology and cross-sectional morphology were observed by SEM. To ensure accuracy, three tests were conducted for each alloy. The dry sliding wear tests were conducted by the block-ring wear tester machine (MRH-3, China) and the test procedure was carried out in accordance with GB/T 12444, as depicted in Fig. 1 (d). The wear specimen had dimensions of 19 mm × 12 mm × 12 mm, and the ring material used as the wear couple was GCr15. The tests were performed under dry conditions at room temperature with a load of 196 N and a rotational speed of 200 rpm for 120 min. To minimize experimental error, three tests were conducted for each alloy. After tests, the three-dimensional microtopography of the worn track and wear volume were determined by a laser confocal scanning microscope (LCSM, OLYMPUS-4100, Japan). As with the abrasive wear test, the worn surface morphology was observed by SEM. The wear rate of the specimen was calculated by the following equation[ 20 ]: $$W=\varDelta V/(F\times D)$$ 1 where W (mm 3 × N − 1 × m − 1 ) is the wear rate, Δ V (mm 3 ) is the wear volume, F (N) is the load and D (m) is the sliding distance. Results and discussions 3.1. Microstructure and phase Table 4 presents the chemical composition of the hardfacing alloys with different Ti contents by OES. The content of C in 0Ti, 3%Ti, and 5%Ti alloys was controlled to be around 1.80 wt. %, whereas the carbon content in 7%Ti alloy showed a slight variation from the rest at 0.40 wt. %. The influence of increasing carbon content on the microstructure will be discussed in detail in the following sections. Table 4 Chemical composition (wt. %) of the hardfacing alloys Alloy C Cr Ti Si Mn Mo V Fe 0Ti 1.82 8.12 0.00 1.59 1.55 1.34 0.56 Bal. 3%Ti 1.76 7.92 2.94 1.63 1.56 1.27 0.52 Bal. 5%Ti 1.77 7.80 5.12 1.56 1.46 1.29 0.56 Bal. 7%Ti 2.23 7.85 6.88 1.60 1.48 1.33 0.53 Bal. Figure 2 presents the XRD results of the four hardfacing alloys. 0Ti alloy was mainly composed of M 7 C 3 phase and austenite phase. The alloys with Ti additions exhibited diffraction peaks of TiC phase and martensite phase. The results indicated the formation of TiC through in-situ reaction of ferrotitanium and graphite in the molten pool. The mass fraction of martensite phase and TiC phase in 3%Ti, 5%Ti, and 7%Ti alloys was calculated by semi-quantitative analysis from the XRD results, with the results presented in Table 5 . As Ti content increased from 0 to 5 wt. %, there was an increase in the intensity of diffraction peaks for both TiC phase and martensite phase, while the intensity of diffraction peaks for M 7 C 3 phase decreased. This suggested the formation of in-situ TiC which led to a decrease in the mass fraction of M 7 C 3 phase and facilitated the development of the martensitic matrix[ 21 ]. As Ti content increased from 5 to 7 wt.%, the mass fraction of TiC phase increased while that of martensite phase decreased. The variation of phases would exert an influence on the properties of the alloys. Table 5 The mass fractions (wt. %) of in-situ TiC and martensite phases Alloy TiC Martensite 0Ti – 0.00 3%Ti 3.52 23.77 5%Ti 5.33 69.97 7%Ti 6.95 51.97 Figure 3 reveals SEM morphology of the four alloys. According to Fig. 3 (a), 0Ti alloy had a matrix structure consisting of primary austenite dendrites, eutectic austenite, and eutectic carbides M 7 C 3 . The lamellar eutectic carbides M 7 C 3 distributed along the grain boundaries of the primary austenite and exhibited a grid-like structure. Figure 3 (b) displays the microstructure of 3%Ti alloy, comprising austenite, elongated eutectic carbides M 7 C 3 , lath martensite, and in-situ TiC particles. Faceted TiC particles were distributed both within the grains and along grain boundaries. The formation of in-situ TiC particles consumed the carbon content in the matrix and promoted the transformation from austenite to martensite[ 22 ]. Lath martensite was formed around the in-situ TiC particles distributed inside the austenitic grains. The distribution of martensite was uneven due to the irregular dispersion of the in-situ TiC particles. In Fig. 3 (c) and Fig. 3 (d), both 5%Ti and 7%Ti alloys consisted of lath martensite, eutectic austenite, eutectic carbides M 7 C 3 , and in-situ TiC particles. However, from 5%Ti to 7%Ti alloy, the morphology of the TiC particles changed from clustered to dendritic. Compared with 3%Ti alloy, refined and dense martensite phase was formed based on the greater transformation from austenite to martensite in 5%Ti and 7%Ti alloy. The phase distribution on the grain boundaries with EDS mapping for 0Ti alloy is depicted in Fig. 4 (a). The EDS results confirmed the presence of minor Mo-C compound and ultrafine martensite around eutectic carbides M 7 C 3 . The compound was granulated with a size of 3 µm. Seungchan Cho et al.[ 23 ] once reported that the addition of Mo to TiC-reinforced Fe composites resulted in the precipitation of MoC-like interfacial phase at the carbides and matrix interfaces. Mo atoms diffused into the interface between eutectic carbides M 7 C 3 and primary austenite and reacted with C atoms to form Mo-C compound. The carbon content was reduced in local microzones around the eutectic carbides during such process. In addition to the formation of Mo-C compound, ultrafine martensite was also formed in 0Ti alloy. The microstructure of the eutectic carbides M 7 C 3 in 5%Ti alloy is shown in Fig. 4 (b). The eutectic carbides M 7 C 3 exhibited a refined and nodular structure, completely replacing the grid-like eutectic carbides M 7 C 3 in 0Ti alloy. Besides, there was also MoC-like interfacial phase formed at the interface between carbides and the matrix in 5%Ti alloy. Based on the analysis of the microstructure, the solidification process of the alloys is illustrated. Figure 5 (a) presents the solidification process of 0Ti alloy. During the supercooling process of liquid phase alloy, solidification initiated with the formation of primary austenite dendrites. When the remaining liquid phase alloy reached the eutectic composition, a mixture of eutectic austenite and eutectic carbides M 7 C 3 formed[ 19 ]. During the growth process of the primary austenite dendrites, because the solid solubility of Cr, C, and Mo elements in austenite was restricted[ 24 ], these alloy elements were then released into the surrounding liquid. Ultimately, lamellar eutectic carbides M 7 C 3 formed along grain boundaries. The precipitated lamellar eutectic carbides M 7 C 3 on the primary austenite grain boundaries formed a grid-like structure, which could enhance the microhardness. Figure 5 (b), (c), and (d) display the solidification process for in-situ TiC-reinforced iron-based hardfacing alloys. The formation of in-situ TiC particles occurred earlier than primary austenite in the molten pool due to its low Gibbs free energy[ 25 ]. In 3%Ti alloy, there were minor fine faceted TiC particles formed primarily in the liquid. When Ti content increased to 5 wt. %, there were more Ti atoms reacting with C atoms to form TiC particles, resulting in an increased amount and large size of TiC particles. Due to the lower density of TiC compared with iron-based alloys, the TiC particles floated in the molten pool[ 26 ]. These floating TiC particles collided, sintered, and finally formed clustered TiC particles in 5%Ti alloy[ 27 ]. With more TiC particles in 7%Ti alloy, the clustered TiC particles aggregated and sintered together to form dendrite TiC particles at the top of the hardfacing layer[ 28 ]. With the low lattice mismatch between TiC and austenite, the in-situ TiC particles could act as nucleation cores for primary austenite[ 29 , 30 ]. Thus, with more nucleation cores, the primary austenite dendrites were refined. In addition, the formation of in-situ TiC consumed C in the liquid, reducing the mass fraction of eutectic carbides M 7 C 3 . At the same time, in-situ TiC particles generated at the grain boundaries could act as nucleation cores for the eutectic phase[ 30 ], resulting in the refinement of eutectic carbides M 7 C 3 , which transformed from a grid-like structure to a nodular structure at the austenite grain boundaries. Due to the decrease in mass fraction of eutectic carbides, more Cr, Mo was solidly dissolved in austenite. The Cr and Mo in austenite could enhanced the hardenability and facilitated the transformation from austenite to martensite[ 31 ]. Despite the same content of elements (Cr, Mn, Mo and V) that could potentially form carbides by combining with carbon, the carbon content in the austenite decreased with increasing Ti content from 0Ti alloy to 5%Ti alloy. The low carbon content in the austenite increased M s , enabling a higher mass fraction of martensite within the matrix. 7%Ti alloy had the highest Ti content among the four alloys, but its martensite mass fraction was lower than that of 5%Ti alloy, which may be attributed to the higher carbon content in 7%Ti alloy (shown in Table 4 ). The mass fraction of martensite would have a significant influence on the properties of the alloys. 3.2. Abrasive wear test Figure 6 illustrates the microhardness and weight loss for four different hardfacing alloys. The microhardness of 0Ti alloy was 416.4 HV 0.2 , slightly lower than that of 3%Ti alloy with 454.3 HV 0.2 . 5%Ti alloy exhibited the highest microhardness of 836.3 HV 0.2 , and 7%Ti alloy had a lower value of 642.1 HV 0.2 . From the weight loss results, 0Ti alloy showed approximately seven times greater weight loss of 3.9989 g compared with 3%Ti alloy of 0.5893 g. 5%Ti alloy exhibited the lowest weight loss of 0.3905 g, whereas 7%Ti alloy had a weight loss of 0.4384 g. Compared with 0Ti alloy, the in-situ TiC-reinforced low chromium iron-based hardfacing alloys demonstrated exceptional wear resistance. The wear resistance was influenced not only by the amount of in-situ TiC particles but also by the phase structure and the phase mass fraction of the matrix[ 32 ]. Among the four different alloys, 5%Ti alloy possessed the highest mass fraction of martensite and an adequate amount of in-situ TiC particles, enabling it to exhibit superior microhardness and wear resistance. Despite the highest mass fraction of TiC in 7%Ti alloy, the increased content of carbon in the matrix led to a decrease in the mass fraction of martensite, resulting in the reduced microhardness and wear resistance. Figure 7 presents the SEM images depicting the morphology of the worn surface of the four different hardfacing alloys after abrasive wear tests. In 0Ti alloy, there were a lot of grooves formed parallel to the wear direction on the austenite matrix, indicating a low wear resistance and a resultant mass loss due to the destruction of the material surface. Meanwhile, with the impact of the abrasive grains, the eutectic carbides M 7 C 3 experienced the fracture and spalled from the matrix, resulting in pits of various diameters. The white particles at the interface between austenite and carbides in Fig. 7 (b) were detected to be the compound of Mo and C. Ref. [ 33 ] found that the existence of the compound of Mo and C at the TiC/austenite interface in TiC-reinforced Fe composites could increase the bonding strength of the interface. As a result, these particles had good adhesion to the matrix and were not removed by the cutting of abrasive grains. In 3%Ti alloy, the fine black particles distributed on the matrix were in-situ TiC particles. Fewer grooves were observed on the worn surface because there were martensite and in-situ TiC particles in the matrix, resulting in relatively higher microhardness and better wear resistance compared with 0Ti alloy. Moreover, the mass fraction of eutectic carbides M 7 C 3 decreased and the pits caused by the spallation of eutectic carbides M 7 C 3 were absent. In Fig. 7 (d), crushed quartz sand was embedded at the center of the groove. In 5%Ti alloy, there were many scratches instead of grooves formed on the worn surface. When abrasive grains were forced to slide on the worn surface of 5%Ti alloy with more martensite and in-situ TiC particles, they could only scratch the surface by microcutting causing the formation of scratches instead of grooves which could remove a lot of material, as shown in Fig. 7 (f). Therefore, the weight loss of 5%Ti alloy was the lowest among the four alloys and exhibited the best wear resistance. In 7%Ti alloy, there were both grooves and scratches on the worn surface due to a decrease in the mass fraction of martensite. As depicted in Fig. 7 (h), the short grooves originated from the matrix and ended where TiC particles were aggregated. However, the compression of the rubber wheel and abrasive grains at the edge of the grooves generated high-stress concentration, resulting in the cracking, brittle fracture, and spallation of TiC particles. Figure 8 presents the SEM images for the cross-sectional worn morphology of the alloys. In 0Ti alloy, the worn surface was uneven with the fracture of eutectic carbides M 7 C 3 and peeling pits caused by the removal of austenite. The primary austenite around the eutectic carbides was removed by the abrasive grains, while eutectic carbides were harder than abrasive grains and would not be scratched by them. After the removal of the austenite matrix, the eutectic carbides were exposed to the surface and then fractured under the impact and abrasion of the abrasive grains[ 34 ]. In 3%Ti alloy, the worn surface was more even than 0Ti alloy. The austenite exhibited poor resistance to the cutting of abrasive grains, resulting in peeling pits. Conversely, the lath martensite effectively resisted the microploughing of abrasive grains, suffering from plastic deformation along the wear direction without being plowed out of the surface substantially. In 5%Ti and 7%Ti alloys, the worn surfaces were nearly flat because there were more mass fraction of martensite and TiC particles. However, microcracks were observed along the grain boundaries. The reason was that eutectic carbides M 7 C 3 precipitated as elongated strips and free nodules at the primary austenite grain boundaries, exhibiting limited plastic deformation capability[ 21 ], whereas the martensite and residual austenite within grains exhibited considerable plastic deformation capability. Therefore, there was inconsistent plastic deformation in the grains surrounding the eutectic carbides M 7 C 3 [ 35 ], which were prone to be the source of microcracks. The microcracks only propagated along the grain boundaries since it was challenging for microcracks to propagate into the austenite and martensite matrix. 3.3. Dry sliding wear test The friction coefficient curves of the dry sliding wear tests for four alloys, along with the average friction coefficient and wear rate, are shown in Fig. 9 . 0Ti alloy had the highest friction coefficient of 0.63, while 5%Ti alloy had the lowest friction coefficient of 0.47. 3%Ti alloy had a friction coefficient similar to that of 7%Ti alloy, with values of 0.56 and 0.54, respectively. It was worth noting that 3%Ti alloy exhibited the highest fluctuation in friction coefficient. The friction coefficient initially increased and then decreased during the first 30 min, which was consistent with the fluctuation commonly observed in most materials during the early stage of the dry sliding wear tests. After the first 30 min, the friction coefficient continued to increase until it reached a stabilized stage after 60 minutes. The fluctuation of the friction coefficient may be attributed to the spallation of TiC particles during wear tests[ 20 ]. The three-dimensional microtopography of the worn track in micro-areas was obtained to calculate the depth of the worn scar and results are shown in Fig. 10 . The selected worn micro-areas were all located in the center of the wear track, which provided the most accurate information of the wear characteristics. As shown in Fig. 10 , the widths of the worn scar were same for four alloys, but the depths of the worn scar varied, indicating different degrees of wear resistance. 0Ti alloy and 3%Ti alloy exhibited a rough boundary between the unworn zone and the worn zone, suggesting a high removal rate of the material. 0Ti alloy exhibited a worn scar with a depth of 70.844 µm and a wear rate of 3.539 × 10 − 6 mm 3 × N − 1 × m − 1 , while 3%Ti alloy had a lower worn scar depth of 60.018 µm and a wear rate of 3.004 × 10 − 6 mm 3 × N − 1 × m − 1 . Both 5%Ti and 7%Ti alloys exhibited a smooth boundary between the unworn zone and the worn zone. 5%Ti alloy demonstrated a worn scar depth of 54.213 µm and a wear rate of 2.117 × 10 − 6 mm 3 × N − 1 × m − 1 and 7%Ti alloy had a slightly higher worn scar depth of 58.307 µm and a wear rate of 2.676 × 10 − 6 mm 3 × N − 1 × m − 1 . Because the austenite was soft, it was easy to be removed from the worn surface by the wear ring. Thus, the wear rate of 0Ti alloy was the highest. As Ti content increased, the mass fraction of martensite and the strength of the matrix was increased. The wear resistance of in-situ TiC-reinforced iron-based hardfacing alloys got improved and the wear rate was decreased. The worn surface morphology of the alloys and the corresponding EDS mapping are shown in Fig. 11 . The worn surface of the four alloys exhibited comparable characteristics, including grooves and tribochemical reaction layers. Compared with 0Ti alloy, the worn surface of 3%Ti, 5%Ti, and 7%Ti alloys exhibited more grooves. During the relative sliding process between the specimen and the wear ring, TiC particles were released onto the friction interface and slid on the surface of the alloys, resulting in the formation of grooves parallel to the sliding direction on the worn surface. Because of the relatively lower mass fraction of martensite in 3%Ti alloy, more matrix was removed from the worn surface and more TiC particles entered the friction interface. The spallation of TiC particles formed pits and the sliding of TiC particles led to grooves. Therefore, the worn surface exhibited more grooves and pits in 3%Ti alloy and the friction curve of 3%Ti alloy showed more fluctuations. In 5%Ti and 7%Ti alloys, the increase in the mass fraction of martensite resulted in stronger strength of the matrix, which prevented the spallation of TiC particles. As a result, there was fewer grooves and pits on the worn surface compared with 3%Ti alloy. As dry sliding wear proceeded, the friction between the specimen and the wear ring caused a high contact surface temperature, leading to tribooxidation at elevated temperature, one of the submechanisms of tribochemical reaction[ 36 ]. In Fig. 11 (a), the O and Cr elements exhibited opposite distribution in the EDS mapping, revealing that austenite grains were the preferential tribooxidation sites. Thus, most of austenite were removed from the worn surface and some austenite grains could directly contact with the wear ring and experienced relative sliding, leading to tribooxidation and the formation of a Fe 2 O 3 oxide layer [ 21 ]. With dry sliding wear progressing, the size of the tribochemical reaction layers enlarged, with its thickness gradually increasing. Compared with 0Ti alloy, the in-situ TiC-reinforced iron-based hardfacing alloy possessed less austenite phase region, so they had tribochemical reaction layers with smaller size. Abrasion during the wear process destroyed a portion of the tribochemical reaction layers, which were then ejected from the contact surface. The remained tribochemical reaction layers were attached to the worn surface and functioned to protect the worn surface[ 37 ]. As a result, the friction coefficient kept stable and decreased slightly in the later stage of the dry sliding wear tests. 3.4. Wear mechanism Figure 12 displays the schematic diagrams of the abrasive wear and dry sliding wear mechanism for both 0Ti alloy and in-situ TiC-reinforced low chromium iron-based hardfacing alloys. During abrasive wear, the abrasive grains continuously slid on the surface of the specimen, ploughing the austenite out of the worn surface, forming pits in 0Ti alloy. The eutectic carbides M 7 C 3 could not be adequately supported by the surrounding austenite, leading to brittle fracture. However, in the in-situ TiC-reinforced low chromium iron-based hardfacing alloys, the martensite provided adequate support to TiC particles, which hindered the sliding of the abrasive grains. Consequently, the worn surface exhibited fewer pits, with less material removal and lower weight loss. To sum up, for 0Ti alloy, the wear mechanism was microploughing, whereas for in-situ TiC-reinforced low chromium iron-based hardfacing alloys, the wear mechanism was microploughing and microcutting. The difference between the dry sliding wear and the abrasive wear was tribochemical reaction layers. During dry sliding wear, the contact surface between the specimen and the wear ring heated up due to friction. The worn surface of 0Ti alloy experienced tribochemical reaction, resulting in the formation of a progressively thicker tribochemical reaction layers. These large size of tribochemical reaction layers were destroyed by abrasion and then ejected from the contact surface, contributing to high friction coefficient and wear rate. However, in in-situ TiC-reinforced low chromium iron-based hardfacing alloys, the presence of martensite and in-situ TiC particles enhanced the strength of the matrix, which was not easy to be removed. Therefore, the small-sized tribochemical reaction layers were formed and adhered firmly to the worn surface. These tribochemical reaction layers could shield the surface from further damage. Therefore, the in-situ TiC reinforced low chromium iron-based hardfacing alloys possessed a lower friction coefficient and wear rate. Conclusions In this study, four kinds of low chromium iron-based hardfacing alloys with Ti contents of 0 wt. %, 3 wt. %, 5 wt. %, and 7 wt. % were prepared by FCAW technology. The effect of Ti content on the microstructure, hardness and wear resistance of low chromium iron-based hardfacing alloys was analyzed, leading to the following general conclusions: The low chromium iron-based hardfacing alloys with different Ti contents consist of martensite, austenite, eutectic carbides M 7 C 3 and in-situ TiC particles. The mass fraction of M 7 C 3 phase and austenite phase declines as Ti content rises from 0 to 5 wt. %, whereas the mass fraction of TiC phase and martensite phase increases. As Ti content further rises to 7 wt. %, the mass fraction of martensite phase decreases, this may be attributed to the increased content of carbon. Among the four different alloys, the alloy with 5 wt. % Ti has the highest hardness because of the highest mass fraction of martensite. The in-situ formation of TiC particles consumes the carbon content in the matrix and raises the martensite start temperature, thus, which facilitates the transformation from austenite to martensite. In the abrasive wear test, with the increase of Ti content, the wear mechanism transforms from microploughing to microcutting. The matrix with more martensite and TiC particles effectively prevents abrasive grains from damaging the surface, and the wear resistance of the alloys is thus improved. In the dry sliding wear test, the main wear mechanism is abrasion and tribochemical reaction. The spalling TiC particles slide on the worn surface to form grooves. However, the tribochemical reaction layers are formed at elevated temperature on the contact surface and protect the worn surface, reducing the friction coefficient and wear rate. Declarations Acknowledgements This work was supported by the Collaborative Innovation Project of the Bureau of Science, Technology and Information of Chaoyang District, Beijing (CYXC2103). 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ZHANG H, QIU F, YANG H Y et al (2022) Microstructure manipulation mechanism and mechanical properties improvement of H13 steel via trace nano-(TiC + TiB2) particles. Materials Characterization 188:111924. https://doi.org/10.1016/j.matchar.2022.111924. KOBAYASHI J, INA D, YOSHIKAWA N et al (2012) Effects of the Addition of Cr, Mo and Ni on the Microstructure and Retained Austenite Characteristics of 0.2% C–Si–Mn–Nb Ultrahigh-strength TRIP-aided Bainitic Ferrite Steels. ISIJ International 52:1894-1901. https://doi.org/10.2355/isijinternational.52.1894. TJONG S C, MA Z Y (2000) Microstructural and mechanical characteristics of in situ metal matrix composites. Materials Science and Engineering: R: Reports 29:49-113. https://doi.org/10.1016/S0927-796X(00)00024-3. LEE J, LEE D, SONG M H et al (2018) In-situ synthesis of TiC/Fe alloy composites with high strength and hardness by reactive sintering. Journal of Materials Science & Technology 34:1397-1404. https://doi.org/10.1016/j.jmst.2017.03.006. A. Fischer (1992) Mechanisms of high temperature sliding abrasion of metallic materials. Wear 152:151-159. https://doi.org/10.1016/0043-1648(92)90210-Y. RAFIEI M, GHAYOUR H, MOSTAAN H et al (2018) The effect of V addition on microstructure and tribological properties of Fe-Ti-C claddings produced by gas tungsten arc welding. Journal of Materials Processing Technology 266:569-578. https://doi.org/10.1016/j.jmatprotec.2018.11.037. Alfons Fischer, Wlodzimierz Dudzinski, Birgit Gleising et al (2018) Analyzing Mild- and Ultra-Mild Sliding Wear of Metallic Materials by Transmission Electron Microscopy. Advanced Analytical Methods in Tribology 2:29-59. https://doi.org/ 10.1007/978-3-319-99897-8.25. Alfons Fischer, Sabine Weiß, Markus A. Wimmer (2012) The tribological difference between biomedical steels and CoCrMo-alloys. Journal of the Mechanical Behavior of Biomedical Materials 9:50-62. https://doi.org/10.1016/j.jmbbm.2012.01.007. Cite Share Download PDF Status: Under Review Version 1 posted Reviewers agreed at journal 02 Nov, 2023 Reviewers invited by journal 02 Nov, 2023 Editor invited by journal 02 Nov, 2023 Editor assigned by journal 31 Oct, 2023 First submitted to journal 30 Oct, 2023 You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-3530208","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Research Article","associatedPublications":[],"authors":[{"id":245697524,"identity":"39660cb3-046b-4921-b243-6dde6e8d6745","order_by":0,"name":"Zhixiang Tong","email":"","orcid":"","institution":"Beijing University of Technology","correspondingAuthor":false,"prefix":"","firstName":"Zhixiang","middleName":"","lastName":"Tong","suffix":""},{"id":245697525,"identity":"e701afe1-9e2b-469e-8d60-293d252e3bae","order_by":1,"name":"Wei Shao","email":"","orcid":"","institution":"Beijing University of 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Technology","correspondingAuthor":true,"prefix":"","firstName":"Dingyong","middleName":"","lastName":"He","suffix":""}],"badges":[],"createdAt":"2023-11-01 02:51:30","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-3530208/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-3530208/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":46045754,"identity":"977cc061-1e91-4b3e-9642-362dd66e11ec","added_by":"auto","created_at":"2023-11-07 22:44:16","extension":"jpg","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":1915256,"visible":true,"origin":"","legend":"\u003cp\u003eSchematic diagrams of (a) the hardfacing processing system, (b) the hardfacing process, (c) the abrasive wear tester, and (d) the dry sliding wear tester\u003c/p\u003e","description":"","filename":"Fig1.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/2e36d56d9fb004a70fafdfb5.jpg"},{"id":46045741,"identity":"dfa23079-65ec-4322-8028-b08f9d0a2fb5","added_by":"auto","created_at":"2023-11-07 22:44:15","extension":"jpg","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":2255112,"visible":true,"origin":"","legend":"\u003cp\u003eX-ray diffraction patterns of the alloys\u003c/p\u003e","description":"","filename":"Fig2.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/9415f673421d805a01264524.jpg"},{"id":46045757,"identity":"b7b4d52c-3f22-42e9-9d38-191fd2eb0931","added_by":"auto","created_at":"2023-11-07 22:44:16","extension":"jpg","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":7580847,"visible":true,"origin":"","legend":"\u003cp\u003eSurface SEM morphology of the alloys: (a) 0Ti, (b) 3%Ti, (c) 5%Ti, (d) 7%Ti\u003c/p\u003e","description":"","filename":"Fig3.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/a96802e6f8060ca489ed9f0d.jpg"},{"id":46046763,"identity":"f79fc9ba-7fd7-4787-9c0d-1cfcc1dcfaac","added_by":"auto","created_at":"2023-11-07 22:52:15","extension":"jpg","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":7563486,"visible":true,"origin":"","legend":"\u003cp\u003eSEM morphology and EDS mapping of the grain boundaries in (a) 0Ti alloy and (b) 5%Ti alloy\u003c/p\u003e","description":"","filename":"Fig4.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/c911d9253e594d75e991b6dc.jpg"},{"id":46045742,"identity":"5d01bbe1-081b-4ed6-a759-a098c107bce0","added_by":"auto","created_at":"2023-11-07 22:44:15","extension":"jpg","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":6865280,"visible":true,"origin":"","legend":"\u003cp\u003eSchematic diagram of the solidification process: (a) 0Ti, (b) 3%Ti, (c) 5%Ti and (d) 7%Ti alloys\u003c/p\u003e","description":"","filename":"Fig5.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/d7b62ba941bb4e7610d3cdc1.jpg"},{"id":46046764,"identity":"084839c2-2661-49fd-bbe9-ba499d2cacae","added_by":"auto","created_at":"2023-11-07 22:52:15","extension":"jpg","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":2726216,"visible":true,"origin":"","legend":"\u003cp\u003eMicrohardness and weight loss of the alloys in the abrasive wear tests\u003c/p\u003e","description":"","filename":"Fig6.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/f1fe5d3c985a964a02866fac.jpg"},{"id":46045746,"identity":"f98af945-66cc-4dd2-a6c3-d19d301e30f1","added_by":"auto","created_at":"2023-11-07 22:44:15","extension":"jpg","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":6797570,"visible":true,"origin":"","legend":"\u003cp\u003eSurface SEM abrasive worn morphology of the alloys: (a, b) 0Ti, (c, d) 3%Ti, (e, f) 5%Ti, (g, h) 7%Ti\u003c/p\u003e","description":"","filename":"Fig7.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/d28566c00da42672aef3d257.jpg"},{"id":46045760,"identity":"3db49ef4-c2fa-4286-9577-c5d94f798a4e","added_by":"auto","created_at":"2023-11-07 22:44:16","extension":"jpg","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":4033492,"visible":true,"origin":"","legend":"\u003cp\u003eCross-sectional SEM worn morphology of the alloys: (a) 0Ti, (b) 3%Ti, (c) 5%Ti and (d) 7%Ti alloys\u003c/p\u003e","description":"","filename":"Fig8.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/bab07c2d16cec31358a38a42.jpg"},{"id":46045752,"identity":"1d6d45b3-8df0-44a9-a7ec-199d8407d884","added_by":"auto","created_at":"2023-11-07 22:44:15","extension":"jpg","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":2691293,"visible":true,"origin":"","legend":"\u003cp\u003eFriction coefficient curves (a) and average friction coefficient and wear rate (b) for alloys\u003c/p\u003e","description":"","filename":"Fig9.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/55ec3f832a06c33209ca5c0f.jpg"},{"id":46046765,"identity":"41a63974-a109-45de-b836-ae337561a597","added_by":"auto","created_at":"2023-11-07 22:52:16","extension":"jpg","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":2603159,"visible":true,"origin":"","legend":"\u003cp\u003eThree-dimensional microtopography of the worn track of alloys: (a) 0Ti, (b) 3%Ti, (c) 5%Ti, (d) 7%Ti, and (e) the worn scar depth\u003c/p\u003e","description":"","filename":"Fig10.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/98aaca2e5369b12af1dfef89.jpg"},{"id":46045751,"identity":"440748d5-1fea-4f83-9cb3-0d69c9512ac3","added_by":"auto","created_at":"2023-11-07 22:44:15","extension":"jpg","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":6934108,"visible":true,"origin":"","legend":"\u003cp\u003eSurface SEM worn morphology with EDS mapping of the alloys: (a) 0Ti, (b) 3%Ti, (c) 5%Ti, (d) 7%Ti alloys\u003c/p\u003e","description":"","filename":"Fig11.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/190ef0baf12f85ab41364cfe.jpg"},{"id":46046766,"identity":"7b736806-0cb3-4127-94dc-6e02cc97f991","added_by":"auto","created_at":"2023-11-07 22:52:16","extension":"jpg","order_by":12,"title":"Figure 12","display":"","copyAsset":false,"role":"figure","size":1685050,"visible":true,"origin":"","legend":"\u003cp\u003eSurface SEM worn morphology with EDS mapping of the alloys: (a) 0Ti, (b) 3%Ti, (c) 5%Ti, (d) 7%Ti alloys\u003c/p\u003e","description":"","filename":"Fig12.tiff.jpg","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/f86c3583fa62cef559eef010.jpg"},{"id":46048141,"identity":"66fdc149-da67-495d-9cf2-47e0a53bdcd5","added_by":"auto","created_at":"2023-11-07 23:00:17","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":2928368,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-3530208/v1/00a8abc7-c715-40c7-ae66-2c1cd47780ce.pdf"}],"financialInterests":"","formattedTitle":"Microstructure and wear resistance of in-situ TiC-reinforced low chromium iron-based hardfacing alloys","fulltext":[{"header":"Introduction","content":"\u003cp\u003eHardbanding is essential for protecting drill pipe joints against abrasion from casing and rock particles. As the environments of oil and gas extraction become increasingly severe, the requirements for hardbanding materials also escalate. [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. The development of hardbanding materials has transformed from high hardness and wear resistance to \"casing-friendly\" materials that possess good wear resistance, low friction coefficient, and anti-spalling properties[\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. One important deposition technique for hardbanding is Flux-Cored Arc Welding (FCAW). Extensive use of FCAW has been found in industries such as aerospace, coal, oilfield, and agriculture due to its high deposition rate, high productivity, simplicity of equipment, and field processability[\u003cspan additionalcitationids=\"CR6\" citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e]. Therefore, the development of \"casing-friendly\" flux-cored wires for FCAW offers an economical solution to address the wear and cost challenges in hardbanding.\u003c/p\u003e \u003cp\u003eHigh chromium white cast irons (12\u0026ndash;30 wt. % Cr, 2.0\u0026ndash;3.6 wt. % C) are an economic and traditional hardbanding material known for high hardness and excellent wear resistance[\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan additionalcitationids=\"CR9 CR10\" citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e]. However, hypereutectic white cast irons with coarse primary carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e (M is generally Cr, Fe and a few other alloy elements) are susceptible to cracking under impact conditions[\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e], and hypoeutectic white cast irons exhibit brittle fracture of carbides under high-stress wear conditions[\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e]. These cracking phenomena greatly limit the service life of hardbanding. Additionally, these materials have a high friction coefficient, which can cause significant casing wear. On the other hand, low chromium iron-based alloys (Cr\u0026thinsp;\u0026lt;\u0026thinsp;12 wt. %, C\u0026thinsp;\u0026lt;\u0026thinsp;2.0 wt. %) are known for excellent fracture strength and resistance to thermal fatigue cracking[\u003cspan additionalcitationids=\"CR13\" citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e]. However, low hardness and low wear resistance limit their suitability for hardbanding applications. To meet the requirements for hardness, wear resistance, and anti-cracking, a new research direction for \"casing-friendly\" hardbanding involves the use of multi-alloy elements to reinforce low chromium iron-based alloys by introducing Ti, Nb, B and other alloy elements into original powder design. As a result, carbides such as TiC, NbC, and B\u003csub\u003e4\u003c/sub\u003eC are precipitated into the weld pool[\u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e, \u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e]. These carbides can act as heterogeneous nucleation cores, effectively refining the microstructure. The hardfacing alloys with refined microstructure possess high hardness and strength. Simultaneously, the carbides dispersed on the surface of the hardfacing alloy can significantly improve wear resistance. Among the carbides, TiC is commonly used as the reinforcing phase in metal matrix composites, including iron-based, nickel-based, aluminum-based, cobalt-based, and titanium-based composites[\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e, \u003cspan additionalcitationids=\"CR17\" citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e]. Shi-Li Shu et al.[\u003cspan citationid=\"CR14\" class=\"CitationRef\"\u003e14\u003c/span\u003e] incorporated TiC nanoparticles into high-Cr hot work die steel, and they found nano-TiC provided nucleation sites for γ-Fe dendrites and prevented γ-Fe growth during solidification. As a result, the steel with TiC nanoparticles exhibited higher strength and toughness. Yan Liu et al.[\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e] prepared aluminum-based composite coatings reinforced with TiC particles using direct laser deposition. They found that the uniformly distributed TiC particles effectively reduced porosity defects and stopped the abrasive grains from pressing into the coating, thus enhancing the wear resistance of the coating. A. Bedolla-Jacuinde et al.[\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e] found that Ti could be used to increase the hardness without affecting fracture toughness by decreasing the volume fraction of eutectic carbides in the 16 wt. % Cr, 2.5 wt. % C white cast iron. However, to our best knowledge, there is limited research on the influence of Ti content on the microstructure and mechanical properties of low chromium iron-based hardfacing alloys.\u003c/p\u003e \u003cp\u003eIn this paper, four kinds of low chromium iron-based hardfacing alloys with different Ti contents were prepared by FCAW technology, and the effect of Ti content on the microstructure, hardness, and wear performance under the abrasive wear and dry sliding wear was analyzed.\u003c/p\u003e"},{"header":"Experimental process","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003e2.1. Material preparation\u003c/h2\u003e \u003cp\u003eMetal-cored flux-cored wires were developed by encasing flux-cored powder with cold-rolled annealed carbon steel strip. The strip, with 0.5 mm thickness and 12 mm width, was shaped into a U-shape, infused with 200\u0026ndash;400 mesh flux-cored powder, and then compressed to a circle with a diameter of 1.6 mm. The powder filling coefficient ranged from 24\u0026ndash;28%. The composition of the flux-cored powder is outlined in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e. Different Ti contents (0, 3, 5 and 7 wt. %) were regulated by the addition of different ferrotitanium contents (0, 16, 26 and 34 wt. %) respectively, as shown in Table\u0026nbsp;\u003cspan refid=\"Tab1\" class=\"InternalRef\"\u003e1\u003c/span\u003e. The obtained alloys are referred to as 0Ti, 3%Ti, 5%Ti and 7%Ti alloys in the paper.\u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab1\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 1\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003eChemical composition (wt. %) of the metal-cored flux-cored wires\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"2\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\"\u003e \u003cp\u003eRaw Powder\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eContent (wt. %)\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eFerrotitanium (70 wt. % Ti)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e0, 16, 26, 34.\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eHigh carbon ferrochrome (68 wt. % Cr, 8 wt. % C)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e40\u0026ndash;50\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eFerromolybdenum (60 wt. % Mo)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e5\u0026ndash;10\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eFerrovanadium (50 wt. % V)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e2\u0026ndash;6\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eElectrolytic manganese (pure)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e6\u0026ndash;10\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eGraphite (pure)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e2\u0026ndash;6\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eFerrosilicon (75 wt. % Si)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e8\u0026ndash;12\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eAluminum (pure)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e0\u0026ndash;4\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eIron (pure)\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003eBal.\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003eThe hardfacing process was conducted on a Q235 steel plate measuring 150 mm \u0026times; 100 mm \u0026times; 100 mm, utilizing a Fronius TPS5000 digital multifunctional welder operated by an ABB-IR60 six-axis robot, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e (a). The chemical composition of the substrate is outlined in Table\u0026nbsp;\u003cspan refid=\"Tab2\" class=\"InternalRef\"\u003e2\u003c/span\u003e. Before hardfacing, the substrate underwent mechanical polishing and cleaning in ethanol. Figure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e (b) illustrates the schematic diagram of the hardfacing process, with the interlayer temperature maintained between 200\u0026deg;C and 300\u0026deg;C. To minimize dilution effects, three layers were deposited and the corresponding hardfacing parameters are provided in Table\u0026nbsp;\u003cspan refid=\"Tab3\" class=\"InternalRef\"\u003e3\u003c/span\u003e.\u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab2\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 2\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003eChemical composition (wt. %) of Q235 steel plate\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"6\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c6\" colnum=\"6\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\"\u003e \u003cp\u003eC\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eSi\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c3\"\u003e \u003cp\u003eMn\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c4\"\u003e \u003cp\u003eP\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c5\"\u003e \u003cp\u003eS\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c6\"\u003e \u003cp\u003eFe\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e0.074\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e0.031\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c3\"\u003e \u003cp\u003e0.101\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c4\"\u003e \u003cp\u003e0.011\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c5\"\u003e \u003cp\u003e0.013\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c6\"\u003e \u003cp\u003e99.710\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab3\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 3\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003eThe welding parameters\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"2\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\"\u003e \u003cp\u003eParameter\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eValue\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eArc voltage\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e22\u0026ndash;24 V\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eWelding current\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e280\u0026ndash;300 A\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eElectrode polarity\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003epositive\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eWire feed speed\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e8.0 m/min\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eWeave width\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e20 mm\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eWelding speed\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e3 mm/s\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eStick-out\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e15\u0026ndash;20 mm\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003eShield gas\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e80% Ar\u0026thinsp;+\u0026thinsp;20% CO\u003csub\u003e2\u003c/sub\u003e\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec4\" class=\"Section2\"\u003e \u003ch2\u003e2.2. Material characterization\u003c/h2\u003e \u003cp\u003eMetallographic specimen measuring 10 mm \u0026times; 10 mm \u0026times; 10 mm were obtained from the hardfacing layer. These specimen were ground, polished, and etched by a solution that contains 3 mL of 68% nitric acid, 15 mL of 38% hydrochloric acid, 3 g of ferric chloride, and 50 mL of deionized water. The phase structure was examined via X-ray diffraction (XRD, BRUKER, D8 ADVANCE, German) with Cu-Kα radiation, with scanning angles ranging from 20 to 95\u0026deg; at room temperature. The microstructure and phase compositions were analyzed by a scanning electron microscope (SEM, FEI, QUANTA-650, Netherlands) equipped with an energy dispersive spectroscopy (EDS, OXFORD, X-MAX, Britain). Additionally, optical emission spectrometry (OES, BRUKER, Q4, German) was employed for chemical composition analysis. The microhardness was measured by a digital microhardness tester (HXD-1000A, China) with 1.96 N load applied for 10 s.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec5\" class=\"Section2\"\u003e \u003ch2\u003e2.3. Wear tests\u003c/h2\u003e \u003cp\u003eThe abrasive wear tests were conducted at room temperature by a wet sand rubber wheel tester machine (MLS-225, China) and the test method was conducted according to ASTM G105, as depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e (c). The wear specimen were cut from the top layer of the hardfacing alloys with a dimension of 57 mm \u0026times; 25mm \u0026times; 10 mm. The rubber wheel had a diameter of 176 mm and a hardness of 60 Shore A, rotating at a speed of 240 rpm. The abrasive slurry consisted of 1 kg water mixed with 1.5 kg quartz sand with the size in the range of 210\u0026ndash;380 \u0026micro;m. The wear test lasted 30 min and a load of 100 N was applied. Before the test, a pre-grinding process was performed for 5 min to eliminate the impact of surface roughness. The specimen were cleaned in ethanol by an ultrasonic cleaner for 3 min before and after the tests, and their weight was measured to an accuracy of 1 mg on an electronic balance (Sartorius, BS224S, China). The worn surface morphology and cross-sectional morphology were observed by SEM. To ensure accuracy, three tests were conducted for each alloy.\u003c/p\u003e \u003cp\u003eThe dry sliding wear tests were conducted by the block-ring wear tester machine (MRH-3, China) and the test procedure was carried out in accordance with GB/T 12444, as depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e (d). The wear specimen had dimensions of 19 mm \u0026times; 12 mm \u0026times; 12 mm, and the ring material used as the wear couple was GCr15. The tests were performed under dry conditions at room temperature with a load of 196 N and a rotational speed of 200 rpm for 120 min. To minimize experimental error, three tests were conducted for each alloy. After tests, the three-dimensional microtopography of the worn track and wear volume were determined by a laser confocal scanning microscope (LCSM, OLYMPUS-4100, Japan). As with the abrasive wear test, the worn surface morphology was observed by SEM. The wear rate of the specimen was calculated by the following equation[\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e]:\u003cdiv id=\"Equ1\" class=\"Equation\"\u003e\u003cdiv format=\"TEX\" class=\"mathdisplay\" id=\"FileID_Equ1\" name=\"EquationSource\"\u003e\n$$W=\\varDelta V/(F\\times D)$$\u003c/div\u003e\u003cdiv class=\"EquationNumber\"\u003e1\u003c/div\u003e\u003c/div\u003e\u003c/p\u003e \u003cp\u003ewhere \u003cem\u003eW\u003c/em\u003e (mm\u003csup\u003e3\u003c/sup\u003e \u0026times; N\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e \u0026times; m\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e) is the wear rate, Δ\u003cem\u003eV\u003c/em\u003e (mm\u003csup\u003e3\u003c/sup\u003e) is the wear volume, \u003cem\u003eF\u003c/em\u003e (N) is the load and \u003cem\u003eD\u003c/em\u003e (m) is the sliding distance.\u003c/p\u003e \u003c/div\u003e"},{"header":"Results and discussions","content":"\u003cdiv id=\"Sec7\" class=\"Section2\"\u003e \u003ch2\u003e3.1. Microstructure and phase\u003c/h2\u003e \u003cp\u003eTable\u0026nbsp;\u003cspan refid=\"Tab4\" class=\"InternalRef\"\u003e4\u003c/span\u003e presents the chemical composition of the hardfacing alloys with different Ti contents by OES. The content of C in 0Ti, 3%Ti, and 5%Ti alloys was controlled to be around 1.80 wt. %, whereas the carbon content in 7%Ti alloy showed a slight variation from the rest at 0.40 wt. %. The influence of increasing carbon content on the microstructure will be discussed in detail in the following sections.\u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab4\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 4\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003eChemical composition (wt. %) of the hardfacing alloys\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"9\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c4\" colnum=\"4\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c5\" colnum=\"5\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c6\" colnum=\"6\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c7\" colnum=\"7\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c8\" colnum=\"8\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c9\" colnum=\"9\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\"\u003e \u003cp\u003eAlloy\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eC\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c3\"\u003e \u003cp\u003eCr\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c4\"\u003e \u003cp\u003eTi\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c5\"\u003e \u003cp\u003eSi\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c6\"\u003e \u003cp\u003eMn\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c7\"\u003e \u003cp\u003eMo\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c8\"\u003e \u003cp\u003eV\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c9\"\u003e \u003cp\u003eFe\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e0Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e1.82\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e8.12\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e0.00\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e1.59\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e \u003cp\u003e1.55\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e \u003cp\u003e1.34\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c8\"\u003e \u003cp\u003e0.56\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c9\"\u003e \u003cp\u003eBal.\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e3%Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e1.76\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e7.92\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e2.94\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e1.63\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e \u003cp\u003e1.56\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e \u003cp\u003e1.27\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c8\"\u003e \u003cp\u003e0.52\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c9\"\u003e \u003cp\u003eBal.\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e5%Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e1.77\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e7.80\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e5.12\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e1.56\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e \u003cp\u003e1.46\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e \u003cp\u003e1.29\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c8\"\u003e \u003cp\u003e0.56\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c9\"\u003e \u003cp\u003eBal.\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e7%Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c2\"\u003e \u003cp\u003e2.23\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e7.85\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c4\"\u003e \u003cp\u003e6.88\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c5\"\u003e \u003cp\u003e1.60\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c6\"\u003e \u003cp\u003e1.48\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c7\"\u003e \u003cp\u003e1.33\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c8\"\u003e \u003cp\u003e0.53\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c9\"\u003e \u003cp\u003eBal.\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e presents the XRD results of the four hardfacing alloys. 0Ti alloy was mainly composed of M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e phase and austenite phase. The alloys with Ti additions exhibited diffraction peaks of TiC phase and martensite phase. The results indicated the formation of TiC through in-situ reaction of ferrotitanium and graphite in the molten pool. The mass fraction of martensite phase and TiC phase in 3%Ti, 5%Ti, and 7%Ti alloys was calculated by semi-quantitative analysis from the XRD results, with the results presented in Table\u0026nbsp;\u003cspan refid=\"Tab5\" class=\"InternalRef\"\u003e5\u003c/span\u003e. As Ti content increased from 0 to 5 wt. %, there was an increase in the intensity of diffraction peaks for both TiC phase and martensite phase, while the intensity of diffraction peaks for M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e phase decreased. This suggested the formation of in-situ TiC which led to a decrease in the mass fraction of M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e phase and facilitated the development of the martensitic matrix[\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e]. As Ti content increased from 5 to 7 wt.%, the mass fraction of TiC phase increased while that of martensite phase decreased. The variation of phases would exert an influence on the properties of the alloys.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003e \u003cdiv class=\"gridtable\"\u003e\u003ctable float=\"Yes\" id=\"Tab5\" border=\"1\"\u003e \u003ccaption language=\"En\"\u003e \u003cdiv class=\"CaptionNumber\"\u003eTable 5\u003c/div\u003e \u003cdiv class=\"CaptionContent\"\u003e \u003cp\u003eThe mass fractions (wt. %) of in-situ TiC and martensite phases\u003c/p\u003e \u003c/div\u003e \u003c/caption\u003e \u003ccolgroup cols=\"3\"\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c1\" colnum=\"1\"\u003e\u003c/div\u003e \u003cdiv align=\"left\" class=\"colspec\" colname=\"c2\" colnum=\"2\"\u003e\u003c/div\u003e \u003cdiv align=\"char\" char=\".\" class=\"colspec\" colname=\"c3\" colnum=\"3\"\u003e\u003c/div\u003e \u003cthead\u003e \u003ctr\u003e \u003cth align=\"left\" colname=\"c1\"\u003e \u003cp\u003eAlloy\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c2\"\u003e \u003cp\u003eTiC\u003c/p\u003e \u003c/th\u003e \u003cth align=\"left\" colname=\"c3\"\u003e \u003cp\u003eMartensite\u003c/p\u003e \u003c/th\u003e \u003c/tr\u003e \u003c/thead\u003e \u003ctbody\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e0Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e\u0026ndash;\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e0.00\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e3%Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e3.52\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e23.77\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e5%Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e5.33\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e69.97\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003ctr\u003e \u003ctd align=\"left\" colname=\"c1\"\u003e \u003cp\u003e7%Ti\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"left\" colname=\"c2\"\u003e \u003cp\u003e6.95\u003c/p\u003e \u003c/td\u003e \u003ctd align=\"char\" char=\".\" colname=\"c3\"\u003e \u003cp\u003e51.97\u003c/p\u003e \u003c/td\u003e \u003c/tr\u003e \u003c/tbody\u003e \u003c/colgroup\u003e \u003c/table\u003e\u003c/div\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e reveals SEM morphology of the four alloys. According to Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e (a), 0Ti alloy had a matrix structure consisting of primary austenite dendrites, eutectic austenite, and eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e. The lamellar eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e distributed along the grain boundaries of the primary austenite and exhibited a grid-like structure. Figure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e (b) displays the microstructure of 3%Ti alloy, comprising austenite, elongated eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e, lath martensite, and in-situ TiC particles. Faceted TiC particles were distributed both within the grains and along grain boundaries. The formation of in-situ TiC particles consumed the carbon content in the matrix and promoted the transformation from austenite to martensite[\u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e]. Lath martensite was formed around the in-situ TiC particles distributed inside the austenitic grains. The distribution of martensite was uneven due to the irregular dispersion of the in-situ TiC particles. In Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e (c) and Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e (d), both 5%Ti and 7%Ti alloys consisted of lath martensite, eutectic austenite, eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e, and in-situ TiC particles. However, from 5%Ti to 7%Ti alloy, the morphology of the TiC particles changed from clustered to dendritic. Compared with 3%Ti alloy, refined and dense martensite phase was formed based on the greater transformation from austenite to martensite in 5%Ti and 7%Ti alloy.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe phase distribution on the grain boundaries with EDS mapping for 0Ti alloy is depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e (a). The EDS results confirmed the presence of minor Mo-C compound and ultrafine martensite around eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e. The compound was granulated with a size of 3 \u0026micro;m. Seungchan Cho et al.[\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e] once reported that the addition of Mo to TiC-reinforced Fe composites resulted in the precipitation of MoC-like interfacial phase at the carbides and matrix interfaces. Mo atoms diffused into the interface between eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e and primary austenite and reacted with C atoms to form Mo-C compound. The carbon content was reduced in local microzones around the eutectic carbides during such process. In addition to the formation of Mo-C compound, ultrafine martensite was also formed in 0Ti alloy. The microstructure of the eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e in 5%Ti alloy is shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e (b). The eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e exhibited a refined and nodular structure, completely replacing the grid-like eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e in 0Ti alloy. Besides, there was also MoC-like interfacial phase formed at the interface between carbides and the matrix in 5%Ti alloy.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eBased on the analysis of the microstructure, the solidification process of the alloys is illustrated. Figure\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e (a) presents the solidification process of 0Ti alloy. During the supercooling process of liquid phase alloy, solidification initiated with the formation of primary austenite dendrites. When the remaining liquid phase alloy reached the eutectic composition, a mixture of eutectic austenite and eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e formed[\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e]. During the growth process of the primary austenite dendrites, because the solid solubility of Cr, C, and Mo elements in austenite was restricted[\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e], these alloy elements were then released into the surrounding liquid. Ultimately, lamellar eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e formed along grain boundaries. The precipitated lamellar eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e on the primary austenite grain boundaries formed a grid-like structure, which could enhance the microhardness. Figure\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e (b), (c), and (d) display the solidification process for in-situ TiC-reinforced iron-based hardfacing alloys. The formation of in-situ TiC particles occurred earlier than primary austenite in the molten pool due to its low Gibbs free energy[\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e]. In 3%Ti alloy, there were minor fine faceted TiC particles formed primarily in the liquid. When Ti content increased to 5 wt. %, there were more Ti atoms reacting with C atoms to form TiC particles, resulting in an increased amount and large size of TiC particles. Due to the lower density of TiC compared with iron-based alloys, the TiC particles floated in the molten pool[\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. These floating TiC particles collided, sintered, and finally formed clustered TiC particles in 5%Ti alloy[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. With more TiC particles in 7%Ti alloy, the clustered TiC particles aggregated and sintered together to form dendrite TiC particles at the top of the hardfacing layer[\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eWith the low lattice mismatch between TiC and austenite, the in-situ TiC particles could act as nucleation cores for primary austenite[\u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e, \u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. Thus, with more nucleation cores, the primary austenite dendrites were refined. In addition, the formation of in-situ TiC consumed C in the liquid, reducing the mass fraction of eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e. At the same time, in-situ TiC particles generated at the grain boundaries could act as nucleation cores for the eutectic phase[\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e], resulting in the refinement of eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e, which transformed from a grid-like structure to a nodular structure at the austenite grain boundaries. Due to the decrease in mass fraction of eutectic carbides, more Cr, Mo was solidly dissolved in austenite. The Cr and Mo in austenite could enhanced the hardenability and facilitated the transformation from austenite to martensite[\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e]. Despite the same content of elements (Cr, Mn, Mo and V) that could potentially form carbides by combining with carbon, the carbon content in the austenite decreased with increasing Ti content from 0Ti alloy to 5%Ti alloy. The low carbon content in the austenite increased M\u003csub\u003es\u003c/sub\u003e, enabling a higher mass fraction of martensite within the matrix. 7%Ti alloy had the highest Ti content among the four alloys, but its martensite mass fraction was lower than that of 5%Ti alloy, which may be attributed to the higher carbon content in 7%Ti alloy (shown in Table\u0026nbsp;\u003cspan refid=\"Tab4\" class=\"InternalRef\"\u003e4\u003c/span\u003e). The mass fraction of martensite would have a significant influence on the properties of the alloys.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec8\" class=\"Section2\"\u003e \u003ch2\u003e3.2. Abrasive wear test\u003c/h2\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003e illustrates the microhardness and weight loss for four different hardfacing alloys. The microhardness of 0Ti alloy was 416.4 HV\u003csub\u003e0.2\u003c/sub\u003e, slightly lower than that of 3%Ti alloy with 454.3 HV\u003csub\u003e0.2\u003c/sub\u003e. 5%Ti alloy exhibited the highest microhardness of 836.3 HV\u003csub\u003e0.2\u003c/sub\u003e, and 7%Ti alloy had a lower value of 642.1 HV\u003csub\u003e0.2\u003c/sub\u003e. From the weight loss results, 0Ti alloy showed approximately seven times greater weight loss of 3.9989 g compared with 3%Ti alloy of 0.5893 g. 5%Ti alloy exhibited the lowest weight loss of 0.3905 g, whereas 7%Ti alloy had a weight loss of 0.4384 g. Compared with 0Ti alloy, the in-situ TiC-reinforced low chromium iron-based hardfacing alloys demonstrated exceptional wear resistance. The wear resistance was influenced not only by the amount of in-situ TiC particles but also by the phase structure and the phase mass fraction of the matrix[\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e]. Among the four different alloys, 5%Ti alloy possessed the highest mass fraction of martensite and an adequate amount of in-situ TiC particles, enabling it to exhibit superior microhardness and wear resistance. Despite the highest mass fraction of TiC in 7%Ti alloy, the increased content of carbon in the matrix led to a decrease in the mass fraction of martensite, resulting in the reduced microhardness and wear resistance.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e presents the SEM images depicting the morphology of the worn surface of the four different hardfacing alloys after abrasive wear tests. In 0Ti alloy, there were a lot of grooves formed parallel to the wear direction on the austenite matrix, indicating a low wear resistance and a resultant mass loss due to the destruction of the material surface. Meanwhile, with the impact of the abrasive grains, the eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e experienced the fracture and spalled from the matrix, resulting in pits of various diameters. The white particles at the interface between austenite and carbides in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e (b) were detected to be the compound of Mo and C. Ref. [\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e] found that the existence of the compound of Mo and C at the TiC/austenite interface in TiC-reinforced Fe composites could increase the bonding strength of the interface. As a result, these particles had good adhesion to the matrix and were not removed by the cutting of abrasive grains. In 3%Ti alloy, the fine black particles distributed on the matrix were in-situ TiC particles. Fewer grooves were observed on the worn surface because there were martensite and in-situ TiC particles in the matrix, resulting in relatively higher microhardness and better wear resistance compared with 0Ti alloy. Moreover, the mass fraction of eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e decreased and the pits caused by the spallation of eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e were absent. In Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e (d), crushed quartz sand was embedded at the center of the groove. In 5%Ti alloy, there were many scratches instead of grooves formed on the worn surface. When abrasive grains were forced to slide on the worn surface of 5%Ti alloy with more martensite and in-situ TiC particles, they could only scratch the surface by microcutting causing the formation of scratches instead of grooves which could remove a lot of material, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e (f). Therefore, the weight loss of 5%Ti alloy was the lowest among the four alloys and exhibited the best wear resistance. In 7%Ti alloy, there were both grooves and scratches on the worn surface due to a decrease in the mass fraction of martensite. As depicted in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003e (h), the short grooves originated from the matrix and ended where TiC particles were aggregated. However, the compression of the rubber wheel and abrasive grains at the edge of the grooves generated high-stress concentration, resulting in the cracking, brittle fracture, and spallation of TiC particles.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e presents the SEM images for the cross-sectional worn morphology of the alloys. In 0Ti alloy, the worn surface was uneven with the fracture of eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e and peeling pits caused by the removal of austenite. The primary austenite around the eutectic carbides was removed by the abrasive grains, while eutectic carbides were harder than abrasive grains and would not be scratched by them. After the removal of the austenite matrix, the eutectic carbides were exposed to the surface and then fractured under the impact and abrasion of the abrasive grains[\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e]. In 3%Ti alloy, the worn surface was more even than 0Ti alloy. The austenite exhibited poor resistance to the cutting of abrasive grains, resulting in peeling pits. Conversely, the lath martensite effectively resisted the microploughing of abrasive grains, suffering from plastic deformation along the wear direction without being plowed out of the surface substantially. In 5%Ti and 7%Ti alloys, the worn surfaces were nearly flat because there were more mass fraction of martensite and TiC particles. However, microcracks were observed along the grain boundaries. The reason was that eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e precipitated as elongated strips and free nodules at the primary austenite grain boundaries, exhibiting limited plastic deformation capability[\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e], whereas the martensite and residual austenite within grains exhibited considerable plastic deformation capability. Therefore, there was inconsistent plastic deformation in the grains surrounding the eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e[\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e], which were prone to be the source of microcracks. The microcracks only propagated along the grain boundaries since it was challenging for microcracks to propagate into the austenite and martensite matrix.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec9\" class=\"Section2\"\u003e \u003ch2\u003e3.3. Dry sliding wear test\u003c/h2\u003e \u003cp\u003eThe friction coefficient curves of the dry sliding wear tests for four alloys, along with the average friction coefficient and wear rate, are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e. 0Ti alloy had the highest friction coefficient of 0.63, while 5%Ti alloy had the lowest friction coefficient of 0.47. 3%Ti alloy had a friction coefficient similar to that of 7%Ti alloy, with values of 0.56 and 0.54, respectively. It was worth noting that 3%Ti alloy exhibited the highest fluctuation in friction coefficient. The friction coefficient initially increased and then decreased during the first 30 min, which was consistent with the fluctuation commonly observed in most materials during the early stage of the dry sliding wear tests. After the first 30 min, the friction coefficient continued to increase until it reached a stabilized stage after 60 minutes. The fluctuation of the friction coefficient may be attributed to the spallation of TiC particles during wear tests[\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe three-dimensional microtopography of the worn track in micro-areas was obtained to calculate the depth of the worn scar and results are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e. The selected worn micro-areas were all located in the center of the wear track, which provided the most accurate information of the wear characteristics. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e, the widths of the worn scar were same for four alloys, but the depths of the worn scar varied, indicating different degrees of wear resistance. 0Ti alloy and 3%Ti alloy exhibited a rough boundary between the unworn zone and the worn zone, suggesting a high removal rate of the material. 0Ti alloy exhibited a worn scar with a depth of 70.844 \u0026micro;m and a wear rate of 3.539 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;6\u003c/sup\u003e mm\u003csup\u003e3\u003c/sup\u003e \u0026times; N\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e \u0026times; m\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e, while 3%Ti alloy had a lower worn scar depth of 60.018 \u0026micro;m and a wear rate of 3.004 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;6\u003c/sup\u003e mm \u003csup\u003e3\u003c/sup\u003e \u0026times; N\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e \u0026times; m\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e. Both 5%Ti and 7%Ti alloys exhibited a smooth boundary between the unworn zone and the worn zone. 5%Ti alloy demonstrated a worn scar depth of 54.213 \u0026micro;m and a wear rate of 2.117 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;6\u003c/sup\u003e mm\u003csup\u003e3\u003c/sup\u003e \u0026times; N\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e \u0026times; m\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e and 7%Ti alloy had a slightly higher worn scar depth of 58.307 \u0026micro;m and a wear rate of 2.676 \u0026times; 10\u003csup\u003e\u0026minus;\u0026thinsp;6\u003c/sup\u003e mm \u003csup\u003e3\u003c/sup\u003e \u0026times; N\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e \u0026times; m\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e. Because the austenite was soft, it was easy to be removed from the worn surface by the wear ring. Thus, the wear rate of 0Ti alloy was the highest. As Ti content increased, the mass fraction of martensite and the strength of the matrix was increased. The wear resistance of in-situ TiC-reinforced iron-based hardfacing alloys got improved and the wear rate was decreased.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eThe worn surface morphology of the alloys and the corresponding EDS mapping are shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003e. The worn surface of the four alloys exhibited comparable characteristics, including grooves and tribochemical reaction layers. Compared with 0Ti alloy, the worn surface of 3%Ti, 5%Ti, and 7%Ti alloys exhibited more grooves. During the relative sliding process between the specimen and the wear ring, TiC particles were released onto the friction interface and slid on the surface of the alloys, resulting in the formation of grooves parallel to the sliding direction on the worn surface. Because of the relatively lower mass fraction of martensite in 3%Ti alloy, more matrix was removed from the worn surface and more TiC particles entered the friction interface. The spallation of TiC particles formed pits and the sliding of TiC particles led to grooves. Therefore, the worn surface exhibited more grooves and pits in 3%Ti alloy and the friction curve of 3%Ti alloy showed more fluctuations. In 5%Ti and 7%Ti alloys, the increase in the mass fraction of martensite resulted in stronger strength of the matrix, which prevented the spallation of TiC particles. As a result, there was fewer grooves and pits on the worn surface compared with 3%Ti alloy.\u003c/p\u003e \u003cp\u003eAs dry sliding wear proceeded, the friction between the specimen and the wear ring caused a high contact surface temperature, leading to tribooxidation at elevated temperature, one of the submechanisms of tribochemical reaction[\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e]. In Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003e (a), the O and Cr elements exhibited opposite distribution in the EDS mapping, revealing that austenite grains were the preferential tribooxidation sites. Thus, most of austenite were removed from the worn surface and some austenite grains could directly contact with the wear ring and experienced relative sliding, leading to tribooxidation and the formation of a Fe\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e oxide layer [\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e]. With dry sliding wear progressing, the size of the tribochemical reaction layers enlarged, with its thickness gradually increasing. Compared with 0Ti alloy, the in-situ TiC-reinforced iron-based hardfacing alloy possessed less austenite phase region, so they had tribochemical reaction layers with smaller size. Abrasion during the wear process destroyed a portion of the tribochemical reaction layers, which were then ejected from the contact surface. The remained tribochemical reaction layers were attached to the worn surface and functioned to protect the worn surface[\u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e]. As a result, the friction coefficient kept stable and decreased slightly in the later stage of the dry sliding wear tests.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec10\" class=\"Section2\"\u003e \u003ch2\u003e3.4. Wear mechanism\u003c/h2\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig12\" class=\"InternalRef\"\u003e12\u003c/span\u003e displays the schematic diagrams of the abrasive wear and dry sliding wear mechanism for both 0Ti alloy and in-situ TiC-reinforced low chromium iron-based hardfacing alloys. During abrasive wear, the abrasive grains continuously slid on the surface of the specimen, ploughing the austenite out of the worn surface, forming pits in 0Ti alloy. The eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e could not be adequately supported by the surrounding austenite, leading to brittle fracture. However, in the in-situ TiC-reinforced low chromium iron-based hardfacing alloys, the martensite provided adequate support to TiC particles, which hindered the sliding of the abrasive grains. Consequently, the worn surface exhibited fewer pits, with less material removal and lower weight loss. To sum up, for 0Ti alloy, the wear mechanism was microploughing, whereas for in-situ TiC-reinforced low chromium iron-based hardfacing alloys, the wear mechanism was microploughing and microcutting.\u003c/p\u003e \u003cp\u003eThe difference between the dry sliding wear and the abrasive wear was tribochemical reaction layers. During dry sliding wear, the contact surface between the specimen and the wear ring heated up due to friction. The worn surface of 0Ti alloy experienced tribochemical reaction, resulting in the formation of a progressively thicker tribochemical reaction layers. These large size of tribochemical reaction layers were destroyed by abrasion and then ejected from the contact surface, contributing to high friction coefficient and wear rate. However, in in-situ TiC-reinforced low chromium iron-based hardfacing alloys, the presence of martensite and in-situ TiC particles enhanced the strength of the matrix, which was not easy to be removed. Therefore, the small-sized tribochemical reaction layers were formed and adhered firmly to the worn surface. These tribochemical reaction layers could shield the surface from further damage. Therefore, the in-situ TiC reinforced low chromium iron-based hardfacing alloys possessed a lower friction coefficient and wear rate.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e"},{"header":"Conclusions","content":"\u003cp\u003eIn this study, four kinds of low chromium iron-based hardfacing alloys with Ti contents of 0 wt. %, 3 wt. %, 5 wt. %, and 7 wt. % were prepared by FCAW technology. The effect of Ti content on the microstructure, hardness and wear resistance of low chromium iron-based hardfacing alloys was analyzed, leading to the following general conclusions:\u003c/p\u003e\n\u003col style=\"list-style-type: lower-roman;\"\u003e\n \u003cli\u003eThe low chromium iron-based hardfacing alloys with different Ti contents consist of martensite, austenite, eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e and in-situ TiC particles. The mass fraction of M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e phase and austenite phase declines as Ti content rises from 0 to 5 wt. %, whereas the mass fraction of TiC phase and martensite phase increases. As Ti content further rises to 7 wt. %, the mass fraction of martensite phase decreases, this may be attributed to the increased content of carbon.\u003c/li\u003e\n \u003cli\u003eAmong the four different alloys, the alloy with 5 wt. % Ti has the highest hardness because of the highest mass fraction of martensite. The in-situ formation of TiC particles consumes the carbon content in the matrix and raises the martensite start temperature, thus, which facilitates the transformation from austenite to martensite.\u003c/li\u003e\n \u003cli\u003eIn the abrasive wear test, with the increase of Ti content, the wear mechanism transforms from microploughing to microcutting. The matrix with more martensite and TiC particles effectively prevents abrasive grains from damaging the surface, and the wear resistance of the alloys is thus improved.\u003c/li\u003e\n \u003cli\u003eIn the dry sliding wear test, the main wear mechanism is abrasion and tribochemical reaction. The spalling TiC particles slide on the worn surface to form grooves. However, the tribochemical reaction layers are formed at elevated temperature on the contact surface and protect the worn surface, reducing the friction coefficient and wear rate.\u003c/li\u003e\n\u003c/ol\u003e"},{"header":"Declarations","content":"\u003ch2\u003eAcknowledgements\u003c/h2\u003e\n\u003cp\u003eThis work was supported by the Collaborative Innovation Project of the Bureau of Science, Technology and Information of Chaoyang District, Beijing (CYXC2103).\u003c/p\u003e\n\u003ch3\u003eCompeting Interests\u003c/h3\u003e\n\u003cp\u003eThe authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n\u003cli\u003eTRUHAN J, MENON R, LECLAIRE F et al (2007) The friction and wear of various hard-face claddings for deep-hole drilling. 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Journal of the Mechanical Behavior of Biomedical Materials 9:50-62. https://doi.org/10.1016/j.jmbbm.2012.01.007.\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":true,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"welding-in-the-world","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"witw","sideBox":"Learn more about [Welding in the World](https://www.springer.com/journal/40194)","snPcode":"40194","submissionUrl":"https://www.editorialmanager.com/witw/","title":"Welding in the World","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"em","reportingPortfolio":"Springer Hybrid","inReviewEnabled":true,"inReviewRevisionsEnabled":false},"keywords":"Flux-Cored Arc Welding, In-situ, TiC, Wear resistance, Low chromium iron-based hardfacing alloys","lastPublishedDoi":"10.21203/rs.3.rs-3530208/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-3530208/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eStrengthening low chromium iron-based alloys with multiple alloy elements is a crucial strategy for developing \"casing-friendly\" hardbanding materials. In this paper, four kinds of low chromium iron-based hardfacing alloys with different Ti contents were prepared by Flux-Cored Arc Welding technology and the effect of Ti content on the microstructure, hardness and wear performance under the abrasive wear and dry sliding wear were analyzed. The in-situ TiC-reinforced iron-based hardfacing alloys exhibited microstructure comprising austenite, martensite, eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e and in-situ TiC particles. The in-situ formation of TiC particles consumed the carbon in the alloy and thus raised the martensite start temperature, resulting in a reduction in the mass fraction of austenite and eutectic carbides M\u003csub\u003e7\u003c/sub\u003eC\u003csub\u003e3\u003c/sub\u003e and an increase in the martensite. Among four hardfacing alloys, the alloy with 5 wt. % Ti exhibited the highest hardness (836.3 HV0.2), the lowest weight loss and the lowest wear rate due to its highest martensite mass fraction (69.97 wt. %). The matrix with in-situ TiC particles and martensite effectively resisted the cutting of abrasive grains and the wear mechanism developed from the microploughing in 0Ti alloy to microcutting in in-situ TiC-reinforced iron-based hardfacing alloys in abrasive wear tests. In dry sliding wear tests, TiC particles were released and slid on the surface of the alloys, resulting in the formation of grooves, while the formation of the tribochemical reaction layers contributed to a reduction in the friction coefficient and wear rate. This study provides a theoretical foundation for the development of \"casing-friendly\" hardbanding materials.\u003c/p\u003e","manuscriptTitle":"Microstructure and wear resistance of in-situ TiC-reinforced low chromium iron-based hardfacing alloys","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2023-11-07 22:44:10","doi":"10.21203/rs.3.rs-3530208/v1","editorialEvents":[{"type":"communityComments","content":0},{"type":"reviewerAgreed","content":"","date":"2023-11-02T12:08:13+00:00","index":0,"fulltext":""},{"type":"reviewersInvited","content":"","date":"2023-11-02T12:04:27+00:00","index":"","fulltext":""},{"type":"editorInvited","content":"Welding in the World","date":"2023-11-02T08:17:47+00:00","index":"","fulltext":""},{"type":"editorAssigned","content":"","date":"2023-10-31T23:28:05+00:00","index":"","fulltext":""},{"type":"submitted","content":"Welding in the World","date":"2023-10-31T01:26:21+00:00","index":"","fulltext":""}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"welding-in-the-world","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"witw","sideBox":"Learn more about [Welding in the World](https://www.springer.com/journal/40194)","snPcode":"40194","submissionUrl":"https://www.editorialmanager.com/witw/","title":"Welding in the World","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"em","reportingPortfolio":"Springer Hybrid","inReviewEnabled":true,"inReviewRevisionsEnabled":false}}],"origin":"","ownerIdentity":"c56d5f19-7247-4f52-ba9b-57ecf58b6dec","owner":[],"postedDate":"November 7th, 2023","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"under-review","subjectAreas":[],"tags":[],"updatedAt":"2023-12-05T21:44:06+00:00","versionOfRecord":[],"versionCreatedAt":"2023-11-07 22:44:10","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-3530208","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-3530208","identity":"rs-3530208","version":["v1"]},"buildId":"_2-kVJe1T_tPrBINL-cwx","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}