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Effects of Co and W on the microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent-clad layers
Journal Pre-proof Effects of Co and W on the microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent-clad layers Yuan-Ching Lin, Yu-Yu Liu PII:
S0043-1648(19)31083-X
DOI:
https://doi.org/10.1016/j.wear.2020.203186
Reference:
WEA 203186
To appear in:
Wear
Received Date: 14 July 2019 Revised Date:
4 January 2020
Accepted Date: 7 January 2020
Please cite this article as: Y.-C. Lin, Y.-Y. Liu, Effects of Co and W on the microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent-clad layers, Wear (2020), doi: https:// doi.org/10.1016/j.wear.2020.203186. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.
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Effects of Co and W on the microstructure and wear behavior of
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NiCrAlMoTiFeNbX equimolar multicomponent-clad layers
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Effects of Co and W on the microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent-clad layers Yuan-Ching Lin* and Yu-Yu Liu Department of Mechanical Engineering, National Taiwan University of Science and Technology, 43, Keelung Road, Section 4, Taipei 10607, Taiwan, ROC
Abstract In this study, we used the gas tungsten arc welding (GTAW) process for cladding NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW multicomponent alloys onto the surface of AISI 1020 low-carbon steel. The microstructures and the sliding wear resistance of the cladding layers were characterized. The experimental results indicated that a multiple-carbide particle composed of (Nb,Ti)C with a TiC-rich core and NbC shell, was synthesized in situ in both multicomponent cladding layers. The eutectic phase of Fe0.875Mo0.125 with the body-centered cubic structure (BCC) was present in the NiCrAlMoTiFeNbCo cladding layer, and the reinforcing phase of Fe2W with the hexagonal-closest-packed (HCP) structure was in the NiCrAlMoTiFeNbCW cladding layer. Wear test results revealed that the wear performance of the multicomponent cladding layers can significantly improve the wear resistance of the AISI 1020 low-carbon steel. The wear resistance of the NiCrAlMoTiFeNbW cladding layer exceeded that of the other cladding layer. The improvement in the wear resistance of the NiCrAlMoTiFeNbW cladding layer was attributable to the high hardness of the cladding layer and the Fe2W reinforcement in the cladding layer. The wear test results confirmed that the cladding layer with higher ratios of hardness to Young’s modulus for both the matrix and the strengthening phase exhibited better wear resistance. Keywords: reinforcing phase, cladding, multicomponent alloys, wear
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1. Introduction Low-carbon steel is a low-cost structural steel used in many industrial fields. However, poor wear resistance limits its range of applications. Many investigations have utilized surface modification to improve the wear resistance of low-carbon steel, such that it can be used to produce moving components of machine tools having low fabrication cost and long life cycle. Thermal spraying, plating, surface deposition of films, layer cladding, and weld cladding have been used to improve the wear performance of steel in various applications. High-entropy alloys, which are novel alloy systems comprising multiple principal elements, have recently been developed [1]. The elements in high-entropy alloys are present in approximately equimolar amounts. The high mixing entropy tends to cause high-entropy alloys to have simple FCC or BCC solid solution structures [2], and they easily form nanoprecipitates and amorphous phases [3]. Additionally, due to the large variety of microstructures and excellent mechanical properties [3, 4], high-entropy alloys are highly versatile and have many areas of application. Recent studies on the eutectic high-entropy alloy of CoCrFeNiNbx have shown that this alloy system has the excellent integrated mechanical properties of ductility and strength [5]. Most bulk ingots of high-entropy alloys have been obtained by arc melting and casting processes [1–6], and they are used in the surface coating of metals to improve the tribological properties of the metallic components of machinery, increasing the wear resistance and
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reducing friction. Additionally, the modified surface can improve corrosion resistance, wear resistance, and thermal stability, thus extending the life of the protected components by more than 70% [7, 8]. Although several multicomponent mixed powders that contain iron [9–12], cobalt [13–16], and nickel-based [17–19] alloy systems have been used for cladding on various substrates, these cladding systems always have a particular major element. Cantor pointed out the multicomponent alloys had an unexpected microstructure, i.e. the number of constituent phases in the microstructure is small, even single-phase, such as a solid solution [20]. He also described some underlying principles of multicomponent and high-entropy alloys and gave some examples of these materials [21]. Additionally, equimolar multicomponent systems exhibit very high hardness, corrosion resistance, and good wear performance under as-cast conditions [1, 2, 7]. Although many studies have been proposed on different issues associated with the equimolar multicomponent alloy systems, the surface cladding of equimolar multicomponent alloy systems has seldom been investigated [22]. Therefore, herein we explored the microstructure and tribological behavior of the in-situ-formed cladding layer of equimolar multicomponent alloys. Various surface cladding methods have been developed to improve the metallurgical quality and tribological properties of the cladding layer [23–25]. At present, the laser cladding method is usually employed for surface cladding, and in some studies, cladding using various powders on metal surfaces has been performed by this method, but this process has a high 3
cost of equipment. To overcome this problem, herein we use the gas tungsten arc welding (GTAW) method for surface cladding, which can rapidly provide a thick cladding layer that forms a metallurgical bond with the substrate, as has been widely studied [26–28]. The wear behavior of components causes performance decline and failure of the machinery, such as positioning errors of robots and contact fatigue of gears. After a long period of cutting operations, obvious wear would occur on the contact surfaces of the slide and the guideway, lowering the accuracy of the machine tool because of the increase in Abbé errors [29]. Thus, in this study, multicomponent cladding layers are synthesized from different equimolar mixed powders to improve the wear properties of the mechanical components, in order to increase the life cycle of the sliding components. Two multicomponent cladding layers were synthesized by an in-situ reaction on the surface of AISI 1020 low-carbon steel, using the GTAW method. NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW powders are easily obtained and extensively used in metallurgical engineering; hence, they were chosen as the cladding materials studied here. In addition, the experimental results concerning the microstructure, constituent phases, hardness, and tribological properties of the in-situ-synthesized cladding layers were analyzed to elucidate the tribological performance of the clad specimens.
2. Experimental 4
AISI 1020 steel substrates containing 0.2 wt.% carbon and having a Vickers hardness of 180 Hv0.3 were used. Each substrate was first processed into a block with dimensions of 150 × 22 × 22 mm by a milling machine, and its surface was then ultrasonically cleaned in acetone for 10 min. To prepare the cladding material, equimolar NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW multicomponent powders were individually mixed using an attritor for 3 h. The mixed powders were then blended with an appropriate amount of an aqueous solution of 4 wt.% polyvinyl alcohol to form pieces of semisolid paste, each with a size of 120 × 6 × 2 mm. The semisolid paste was pre-placed on the surface of each AISI 1020 substrate, which was baked for 24 h at 150 °C to completely remove water and increase the adhesion between the cladding material and the substrate. After that, the cladding process was conducted. Table 1 indicates the welding parameters, which were suitable for the surface cladding of equimolar NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW multicomponent powders. After the cladding process, the various cladding layers were cut to suitable sizes for metallurgical
characterization
and
wear testing;
they were cut
by using wire
electrical-discharge machining (WEDM). An image analyzer was used to calculate the dilution rate of the cladding layer according to the cross-section of the cladding. The dimensions of the wear test specimen were controlled at 22 mm × 6 mm × (2 ± 0.01) mm, and an arc of radius 5 mm was fabricated by WEDM at the wear test end of the specimen to obtain the wear test geometry of the line-contact type. The wear test specimens were not subjected to 5
any heat treatment; hence the wear performance was assessed in the as-clad condition. The surface for the wear test was polished to control the surface roughness, making it be consistent in each specimen. The wear test was performed according to ASTM G99 standard [30], pin-on-disk mode with a rotating tribometer. The counterpart (disk) for the wear test was made of hardened AISI 52100 steel with a hardness of 61 HRc (≈ 720 Hv0.3) and a surface roughness controlled to approximately Ra = 0.06 µm. Each wear test under a particular condition was conducted at least four times to confirm the repeatability of the experiments. Table 2 presents the wear test conditions. After wear testing, the sample was examined by a stereo optical microscope used to measure the dimensions of the wear scars for calculating the wear loss of the wear test specimens. Nano-indentation was conducted to identify the micromechanical properties of the constituent phases in the cladding layer. Microhardness testing, x-ray diffraction (XRD), transmission electron microscopy (TEM), and electron probe microanalysis (EPMA) were used to determine the hardness distribution of the cladding layer, crystalline structure of the constituent phases in the cladding layer, microstructure of the cladding layer, and composition of the constituent phases, respectively. The XRD scanning angle, 2θ, was between 20° and 100°, and the scanning speed was 1°/min; the maximum acceleration voltage of TEM was 300 kV, and EPMA was 30 kV. Selected area diffraction (SAD) analysis was used to identify the crystal structure of the reinforcing phases and matrix in the cladding layer. 6
3. Results and discussion 3-1 Microstructural analyses of cladding layers In the multicomponent alloy-cladding process with the GTAW method, input energy per unit length less than the critical value resulted in the cladding layer failing to bond well on the substrate, and the cladding layer spalled. Nevertheless, the high input energy caused a high dilution rate of the cladding layer. Those phenomena might be attributable to the differences in the melting points of the equimolar components and low heat conductivity of the multicomponent alloys. The image analysis results indicated that the dilution rates of the NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW cladding layers were 64% and 51%, respectively. 3-1-1 Metallurgical characterization of the NiCrAlMoTiFeNbCo cladding layer Figure 1 displays the microstructure of the NiCrAlMoTiFeNbCo cladding layer at various positions. The reinforcements included some particulate and vein-shaped precipitates. In the cladding process, the cooling rates varied with the position, causing the microstructure in the cladding layer to vary. During the solidification process, the high cooling rate of the liquid metal resulted in a high degree of supercooling, and the critical nucleus radius was reduced. That caused the number of the nuclei to be increased in the region of the high cooling rate, and the grain size was smaller than that of the low cooling rate. Additionally, the thermal gradient and the solidification rate are the most important factors that determined the
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grain shape and grain size [31]. Figure 1(a) shows the microstructure at the top of the cladding layer contained a larger number of vein-shaped reinforcements than at other positions; this was because the cooling rate there was higher. The reinforcements in the bottom region were sparser because of the dilution effect of the substrate. However, the interface between the cladding layer and the substrate exhibited a good metallurgical bonding, as shown in Fig. 1(c). The cooling rate in the middle region of the cladding layer was lower than that in all other regions, causing the microstructure there to be coarse. Figures 2 and 3 respectively display the results of the EPMA and XRD analyses of the NiCrAlMoTiFeNbCo cladding layer. According to these results, the matrix was α-Fe; the particulate reinforcements were composed of TiC, (Nb,Ti)C, and NbC, forming multiple-carbide particles with TiC-rich cores; and the vein-shaped reinforcements were composed of the eutectic phase Fe0.875Mo0.125. The EPMA results also indicated that Nb element was present in the Fe0.875Mo0.125 region. Moreover, a few particles of Al2O3 were detected in the cladding layer, and the related research indicated that the Al2O3 particles could contribute at the sites to heterogeneous nucleation of TiC [32], as shown in Fig. 2. A previous investigation demonstrated that the Ti/Nb atomic ratio dominated the formation of carbides when the melting pool temperature was below 2000 K. At temperatures below 2000 K, the Gibbs free energy of TiC (∆GTiC) was less than that of NbC [33]. Therefore, TiC particles formed first in the melting pool, reducing the Ti concentration in a microscale zone around the 8
TiC particles and changing the Ti/Nb ratio such that NbC was formed on the surfaces of the TiC particles. Thus, multiple-carbide particles (i.e., TiC particles wrapped in NbC) were formed. (Ti,Nb)C was also formed at the TiC/NbC interface. The EPMA and XRD analysis and SEM microstructural observation results confirmed this finding. Furthermore, the BEI image of the SEM in Fig. 1(a) revealed that the cores of the reinforcing particles were black, but the rest was grey, indirectly verifying that the reinforcing particle was a multiphase carbide. This result may be caused by the high cooling rate during solidification in the cladding process, which was different from the low cooling rate in the casting process [33, 34]. Figure 4 displays the TEM characterization results of the cladding layer. The reinforcing particles were distributed in the matrix, as depicted in Figs. 4(a) and (b). Figures 4(c) and (d) show the results of the selected area diffraction (SAD) analyses of the vein-shaped precipitates and the matrix, respectively, revealing that their crystal structures were BCC, but their lattice constants were 0.3046 nm for Fe0.875Mo0.125 and 0.2845 nm for α-Fe.
3-1-2 Metallurgical characterization of the NiCrAlMoTiFeNbW cladding layer Figure 5 shows the microstructure of the NiCrAlMoTiFeNbW cladding layer at various positions. The reinforcements included some particulates and petal-shaped precipitates. The microstructure at the top of the cladding layer contained more and smaller petal-shaped 9
precipitates compared to that at the other positions because the cooling rate was higher there. The dilution effect of the substrate caused the petal-shaped precipitates at the bottom region to be sparse and small in the NiCrAlMoTiFeNbW cladding layer. A few vein-shaped precipitates were also found in the bottom region of the NiCrAlMoTiFeNbW cladding layer. The interface between the bottom of the cladding layer and the substrate exhibited favorable metallurgical bonds. The cooling rate was the lowest in the middle of the cladding layer, causing the petal-shaped precipitates in this region to be coarse. Figures 6 and 7 show the EPMA and XRD results of the NiCrAlMoTiFeNbW cladding layer, respectively. According to these results, the matrix was α-Fe; the reinforcement comprised of TiC, (Nb,Ti)C, and NbC multiple-carbides; and the petal-shaped precipitates were composed of the Fe2W phase. A few Al2O3 particles were also detected in the cladding layer. The EPMA characterization showed that the Fe2W phase contained some Mo and Nb elements. The particulate reinforcements were composed of multiple carbides, and they had a TiC core and an NbC outer layer. The EPMA and XRD analyses and the SEM microstructural observations verified this finding. In the BEI image of the SEM in Fig. 5(a), the cores of the reinforcing particles were black, but the rest were gray, indicating that the compositions of the core and the outer layer differed. Figure 8 shows the TEM results of the NiCrAlMoTiFeNbW cladding layer. The reinforcing phases were distributed in the matrix, which is depicted in Fig. 8(a). Figures 8(b) 10
and (c) present the results of the SAD analyses of the reinforcement and the matrix, respectively. Figure 8(b) reveals that the crystalline structure of the reinforcement (Fe2W) was HCP with lattice constants of a = 0.5093 nm and c = 0.7470 nm. Figure 8(c) displays that the crystalline structure of the matrix was BCC with a lattice constant of a = 0.2751 nm. 3-2 Mechanical properties of the cladding layers The hardness (H) and Young’s modulus (E) of the cladding layers were characterized using a nanoindenter, and then the ratio of the H/E of each phase was calculated. Table 3 indicates that the hardness of the reinforcement Fe2W in the NiCrAlMoTiFeNbW cladding layer exceeded that of Fe0.875Mo0.125 in the NiCrAlMoTiFeNbCo cladding layer. Nevertheless, the hardness of the matrix of the NiCrAlMoTiFeNbW cladding layer was less than that of the matrix of the NiCrAlMoTiFeNbCo cladding layer. The EPMA analyses demonstrated that the matrix of the NiCrAlMoTiFeNbCo cladding layer contained Ni, Cr, Mo, and Co, while that of the NiCrAlMoTiFeNbW cladding layer contained Ni, Cr, Mo, and a little W. Furthermore, the lattice constant of the matrix of the NiCrAlMoTiFeNbCo cladding layer was greater than that of the matrix of the NiCrAlMoTiFeNbW cladding layer. These results implied that the solid solution strengthening effect of Co in the matrix was stronger than that of W in this cladding process. In contrast, the NiCrAlMoTiFeNbW cladding layer contained more of the reinforcing phases compared to the NiCrAlMoTiFeNbCo cladding layer because of the element W. This phenomenon was also related to the dilution rate of the NiCrAlMoTiFeNbW cladding layer, 11
which was less than that of the NiCrAlMoTiFeNbCo cladding layer. Therefore, the microhardness of the NiCrAlMoTiFeNbW cladding layer exceeded that of the NiCrAlMoTiFeNbCo cladding layer. Figure 9 presents the distribution of the microhardness from the cladding surface to the substrate, at a load of 300 g. 3-3 Tribological behaviors of the cladding layers
Figure 10 shows the wear loss of each specimen at different sliding speeds. The wear test results revealed that the wear resistance of all cladding specimens obviously improved, especially at a high sliding speed, compared to that of the AISI 1020 steel substrate. The wear loss of AISI 1020 steel at the high sliding speed was more than double that at low sliding speed because the former produced a high temperature in the contact region, and this high temperature led to the softening of the AISI 1020 steel along with plastic deformation accompanied by severe wear.
Figure 11 shows a typically worn surface of the AISI 1020 steel at a sliding speed of 1.643 m/s. The worn surface had the characteristics of plastic flow accompanying adhesive wear. Initially, the line contact mode caused high contact stress, higher than the yield stress of the low-carbon steel. Therefore, the real contact stress caused the high-stress asperities to experience plastic deformation, in which the accompanying shear effect led to adhesive wear behavior. After that, the apparent contacting area increased gradually with the wear, and the metal transfer and accumulated debris made scratches on the worn surface during sliding. 12
EDS characterization confirmed that the transfer film contained the elements from the counterpart and some oxides.
Figure 12 displays typical fluctuation of the friction coefficient during the wear test at high sliding speed for all the test specimens. The friction coefficient varied over a range of approximately 0.45–0.65. Moreover, the frictional behaviors of the wear test specimens at the low sliding speed were similar to that of the wear test specimens at the high sliding speed.
Figure 13 presents a typically worn surface of the NiCrAlMoTiFeNbCo-clad specimen, in which some patches of a “glaze” layer were formed. A previous study indicated that it is an oxide film [35]. The effect of precipitation strengthening of the reinforcements and solid solution strengthening reduced the plastic flow and adhesive behavior of the cladding layer that resulted in “glaze” oxide layer formation on the high-temperature rubbing surface and no severe wear on the sliding surface. However, the wear behavior during the sliding process caused the detachment of particulate reinforcements and made some small pits on the worn surface. The positions of the detached particulate reinforcements are clearly showing in the BEI image, Fig. 13(b), as black points. The detached particles contributed as the third bodies on the sliding surface, accompanying oxide film formation to reduce the friction coefficient [36]. Therefore, the friction coefficient of the NiCrAlMoTiFeNbCo-clad specimen was lower than that of the AISI 1020 steel specimen, varying over the range of approximately 0.4–0.6.
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Additionally, the vein-shaped reinforcements improved the wear performance effectively because they can bond strongly on the sliding surface by mechanical locking due to complex geometric characteristics.
The NiCrAlMoTiFeNbW cladding specimen had the best wear performance of any of the wear test specimens, and the friction coefficient of that varied over a range of approximately 0.4–0.58. Figure 14 shows the worn surface of the NiCrAlMoTiFeNbW cladding specimen subjected to a low sliding speed. The worn surface displays petal-shaped precipitates that had a uniform distribution in the worn surface, bonding well with the matrix to resist the penetration of the counterpart and thus to improve the wear performance. Moreover, the thin “glaze” oxide film covered the worn surface completely. Figure 14(b) has a BEI image that indicates these petal-shaped precipitates distributed themselves on the worn surface more clearly, and the EDS semi-quantitative analysis (the diamond mark in Fig. 14 (b)) results indicated that worn surface contained oxygen, confirming there was a “glaze” oxide layer on the worn surface [37]. Mild abrasive wear with oxidation wear characteristics was observed, and the reinforcing phase Fe2W emerged on the worn surface. The Fe2W phase exhibited strong bonding with the matrix; it did not detach from the matrix during sliding, so the wear resistance of the cladding layer was improved effectively. Therefore, at a low sliding speed, the wear loss of the NiCrAlMoTiFeNbW-clad specimen was approximately 25% of that of the NiCrAlMoTiFeNbCo-clad specimen; however, at a high sliding speed, it was approximately 14
76%.
Figure 15 displays the worn surface of the NiCrAlMoTiFeNbW cladding specimen obtained after it was subjected to a high sliding speed. The shallow abrasive wear characteristics were observed, and some cracks were found in the petal-shaped reinforcing phase, which had the highest hardness than the other reinforcement, as shown in Table 3. At a high sliding speed, the temperature at the rubbing surfaces was high, leading to high thermal stress in the reinforcement, caused by the difference of thermal conductivity between matrix and reinforcement. Meanwhile, the traction force promoted the cracks to propagate normal to the sliding direction, resulting in thermal fatigue and detachment. As a result, the wear rate increased at the high sliding speed condition. In contrast, the wear loss of the NiCrAlMoTiFeNbCo cladding specimen at the high sliding speed was less than that at the low sliding speed because at the high sliding speed, the oxide film growth rate was high, which could reduce the adhesive behaviors. In all cladding specimens, the friction coefficient at the high sliding speed was less than that at the low sliding speed because the lower-friction oxide film formed quickly on the rubbing surface at the high temperature.
According to the experimental results, the NiCrAlMoTiFeNbW cladding layer exhibited excellent sliding wear resistance under the dry sliding wear test condition because Fe2W precipitate had good bonding with the matrix in the cladding layer, considerably improving
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the wear resistance of the cladding layer. A previous investigation demonstrated that the ratio of hardness to Young’s modulus (H/E) significantly affects the wear resistance of a nanocomposite coating [38]. High H/E was associated with favorable wear performance. Table 3 shows the H/E values for the major phases in the clad specimens. The H/E ratios of both the matrix and the reinforcing phase Fe2W in the NiCrAlMoTiFeNbW cladding layer were higher than those of both the matrix and the reinforcing phase Fe0.875Mo0.125 in the NiCrAlMoTiFeNbCo cladding layer, respectively. For a fixed specimen hardness, a low Young’s modulus is associated with a high H/E ratio; therefore, at a given contact load, a lower E results in a higher elastic strain of the surface asperities, providing more contact positions and a larger real contact area. This effectively reduced the real contact stress and wear behavior. Accordingly, the NiCrAlMoTiFeNbW cladding layer had the best wear performance among the wear test specimens herein.
4. Conclusion The following conclusions are drawn from the results of the microstructural analysis and the wear testing: 1. The reinforcing particles formed in situ during surface cladding in both the NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW cladding layers, comprised multiple layers of phases (Nb,Ti)C/NbC with a TiC-rich core. 2. The vein-shaped reinforcing phase Fe0.875Mo0.125 in the NiCrAlMoTiFeNbCo cladding layer and the petal-shaped reinforcing phase Fe2W in the NiCrAlMoTiFeNbW cladding layer were formed in situ during the cladding process, thereby increasing the hardness and improving the wear performance of the cladding layers. 16
3. The NiCrAlMoTiFeNbW cladding layer exhibited the best wear performance because this cladding layer had the highest H/E ratio and the Fe2W reinforcement phase was significantly harder than the Fe0.875Mo0.125 reinforcement phase. 4. The matrices of the different cladding layers had the same BCC crystal structure but different lattice constants because of solid solution strengthening containing different elements. 5. The multicomponent cladding process improved the tribological performance of the low-carbon steel. This process can extend the application of low-carbon steel to the moving components of a machine.
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20
Appendix: Wear volume calculation for the pin specimen
θ1 O
θ2 C
The shadow plane represents the worn surface of the pin specimen with a cylindrical end with a radius of curvature R Definition of symbols: L0: the width of the pin specimen R: radius of curvature of the cylindrical end θ1: semi-angle of the center on the small worn side θ2: semi-angle of the center on the large worn side α: semi-angle of the center at an arbitrary position of longitudinal cross section of the worn plane H1: width of the wear on the small worn side H2: width of the wear on the large worn side H: width of the wear of the longitudinal cross section of the worn plane at arbitrary position A: the area of the longitudinal cross section of the worn volume at an arbitrary position V: wear volume of the pin specimen Assume H2 ≧ H1, and θ2 ≦ 90∘ 21
Then,
，
,
If 1 = 2, 1 = 2 = , then the wear volume of the pin specimen can be calculated as follows. assume
，let
，
Then the term in the parentheses of the last term of Equation (A1) can be expressed as follow.
22
Use L'Hospital's Rule for eq. (A4)
∴ when
,
≈0
Note: the angle must be expressed using the unit of radian
23
Table 1 cladding parameter of each specimen Cladding materials
NiCrAlMoTiFeNbCo
NiCrAlMoTiFeNbW
Welding current (A)
110
Welding voltage (V)
17
Traveling speed(cm/min)
6.0
5.5
Flow rate of Ar (L/min)
10
Electrode
W-2% thorium
Polarity of electrode
DCEN (DCSP)
Diameter of W electrode (mm) Gap between electrode and workpiece (mm) Input energy per traveling length (kJ/cm)
1.6 2 18.7
20.4
Table 2. wear specimens and test conditions Dimension of the specimens (mm) Pin /Disk
Elastic modules (GPa)
22
6
2 with R5
186
NiCrAlMoTiFeNbW
186
AISI 52100 steel
210
0.29 / 0.3
(cladding layer / AISI 52100) Load (kg)
3.5
Maximum contact stress (MPa)
343
Sliding distance (m) Temperature Lubrication
11.5
NiCrAlMoTiFeNbCo
Poisson ratio
Sliding speed (m/s)
φ60
0.817
1.634 1470 25 Dry sliding
Table 3. Mechanical properties and the ratio of hardness to Young’s Modulus of each phase in the clad specimens NiCrAlMoTiFeNbCo clad layer Phases
Hardness (GPa) Young’s modulus(GPa)
H/E
Fe0.875Mo0.125 reinforcement
10.46
225
0.0465
Matrix
7.38
197
0.0375
NiCrAlMoTiFeNbW clad layer Fe2W reinforcement
16.34
246
0.0664
Matrix
7.20
186
0.0387
* Young’s modulus and hardness were inspected by nanoindenter * indentation load 2500 µN
Highlights 1. Equimolar multicomponent alloyed in-situ on steel to improve its wear performance. 2. Formation mechanism of the reinforcing phases in the clad layer was explored. 3. The correlation between the wear of the sample and the index of H/E was discussed. 4. Alloy element effect on the wear resistance of the clad layer was discussed.
Conflicts of Interest Statement
Manuscript title: Effect of Co and W on microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent clad layers Ref: WEA_2019_894
The authors whose names are listed immediately below certify that they have NO affiliations with or involvement in any organization or entity with any financial interest (such as honoraria; educational grants; participation in speakers’ bureaus; membership, employment, consultancies, stock ownership, or other equity interest; and expert testimony or patentlicensing arrangements), or non-financial interest (such as personal or professional relationships, affiliations, knowledge or beliefs) in the subject matter or materials discussed in this manuscript.
Author names: Yuan-Ching Lin Yu-Yu Liu
S0043-1648(19)31083-X
DOI:
https://doi.org/10.1016/j.wear.2020.203186
Reference:
WEA 203186
To appear in:
Wear
Received Date: 14 July 2019 Revised Date:
4 January 2020
Accepted Date: 7 January 2020
Please cite this article as: Y.-C. Lin, Y.-Y. Liu, Effects of Co and W on the microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent-clad layers, Wear (2020), doi: https:// doi.org/10.1016/j.wear.2020.203186. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier B.V.
AUTHORSHIP STATEMENT Manuscript title:
Effects of Co and W on the microstructure and wear behavior of
__
NiCrAlMoTiFeNbX equimolar multicomponent-clad layers
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acquisition of data: __Y. Y. Liu___, _____________, _____________; analysis and/or interpretation of data: ________________________, ___________. Category 2 Drafting the manuscript:
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revising the manuscript critically for important intellectual content: _ Y. C. Lin_ __, ____________, _____________, _____________. Category 3 Approval of the version of the manuscript to be published (the names of all authors must be listed): _ Y. C. Lin____, _______________, _______________, _______________, _______________, _______________.
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Effects of Co and W on the microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent-clad layers Yuan-Ching Lin* and Yu-Yu Liu Department of Mechanical Engineering, National Taiwan University of Science and Technology, 43, Keelung Road, Section 4, Taipei 10607, Taiwan, ROC
Abstract In this study, we used the gas tungsten arc welding (GTAW) process for cladding NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW multicomponent alloys onto the surface of AISI 1020 low-carbon steel. The microstructures and the sliding wear resistance of the cladding layers were characterized. The experimental results indicated that a multiple-carbide particle composed of (Nb,Ti)C with a TiC-rich core and NbC shell, was synthesized in situ in both multicomponent cladding layers. The eutectic phase of Fe0.875Mo0.125 with the body-centered cubic structure (BCC) was present in the NiCrAlMoTiFeNbCo cladding layer, and the reinforcing phase of Fe2W with the hexagonal-closest-packed (HCP) structure was in the NiCrAlMoTiFeNbCW cladding layer. Wear test results revealed that the wear performance of the multicomponent cladding layers can significantly improve the wear resistance of the AISI 1020 low-carbon steel. The wear resistance of the NiCrAlMoTiFeNbW cladding layer exceeded that of the other cladding layer. The improvement in the wear resistance of the NiCrAlMoTiFeNbW cladding layer was attributable to the high hardness of the cladding layer and the Fe2W reinforcement in the cladding layer. The wear test results confirmed that the cladding layer with higher ratios of hardness to Young’s modulus for both the matrix and the strengthening phase exhibited better wear resistance. Keywords: reinforcing phase, cladding, multicomponent alloys, wear
1
1. Introduction Low-carbon steel is a low-cost structural steel used in many industrial fields. However, poor wear resistance limits its range of applications. Many investigations have utilized surface modification to improve the wear resistance of low-carbon steel, such that it can be used to produce moving components of machine tools having low fabrication cost and long life cycle. Thermal spraying, plating, surface deposition of films, layer cladding, and weld cladding have been used to improve the wear performance of steel in various applications. High-entropy alloys, which are novel alloy systems comprising multiple principal elements, have recently been developed [1]. The elements in high-entropy alloys are present in approximately equimolar amounts. The high mixing entropy tends to cause high-entropy alloys to have simple FCC or BCC solid solution structures [2], and they easily form nanoprecipitates and amorphous phases [3]. Additionally, due to the large variety of microstructures and excellent mechanical properties [3, 4], high-entropy alloys are highly versatile and have many areas of application. Recent studies on the eutectic high-entropy alloy of CoCrFeNiNbx have shown that this alloy system has the excellent integrated mechanical properties of ductility and strength [5]. Most bulk ingots of high-entropy alloys have been obtained by arc melting and casting processes [1–6], and they are used in the surface coating of metals to improve the tribological properties of the metallic components of machinery, increasing the wear resistance and
2
reducing friction. Additionally, the modified surface can improve corrosion resistance, wear resistance, and thermal stability, thus extending the life of the protected components by more than 70% [7, 8]. Although several multicomponent mixed powders that contain iron [9–12], cobalt [13–16], and nickel-based [17–19] alloy systems have been used for cladding on various substrates, these cladding systems always have a particular major element. Cantor pointed out the multicomponent alloys had an unexpected microstructure, i.e. the number of constituent phases in the microstructure is small, even single-phase, such as a solid solution [20]. He also described some underlying principles of multicomponent and high-entropy alloys and gave some examples of these materials [21]. Additionally, equimolar multicomponent systems exhibit very high hardness, corrosion resistance, and good wear performance under as-cast conditions [1, 2, 7]. Although many studies have been proposed on different issues associated with the equimolar multicomponent alloy systems, the surface cladding of equimolar multicomponent alloy systems has seldom been investigated [22]. Therefore, herein we explored the microstructure and tribological behavior of the in-situ-formed cladding layer of equimolar multicomponent alloys. Various surface cladding methods have been developed to improve the metallurgical quality and tribological properties of the cladding layer [23–25]. At present, the laser cladding method is usually employed for surface cladding, and in some studies, cladding using various powders on metal surfaces has been performed by this method, but this process has a high 3
cost of equipment. To overcome this problem, herein we use the gas tungsten arc welding (GTAW) method for surface cladding, which can rapidly provide a thick cladding layer that forms a metallurgical bond with the substrate, as has been widely studied [26–28]. The wear behavior of components causes performance decline and failure of the machinery, such as positioning errors of robots and contact fatigue of gears. After a long period of cutting operations, obvious wear would occur on the contact surfaces of the slide and the guideway, lowering the accuracy of the machine tool because of the increase in Abbé errors [29]. Thus, in this study, multicomponent cladding layers are synthesized from different equimolar mixed powders to improve the wear properties of the mechanical components, in order to increase the life cycle of the sliding components. Two multicomponent cladding layers were synthesized by an in-situ reaction on the surface of AISI 1020 low-carbon steel, using the GTAW method. NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW powders are easily obtained and extensively used in metallurgical engineering; hence, they were chosen as the cladding materials studied here. In addition, the experimental results concerning the microstructure, constituent phases, hardness, and tribological properties of the in-situ-synthesized cladding layers were analyzed to elucidate the tribological performance of the clad specimens.
2. Experimental 4
AISI 1020 steel substrates containing 0.2 wt.% carbon and having a Vickers hardness of 180 Hv0.3 were used. Each substrate was first processed into a block with dimensions of 150 × 22 × 22 mm by a milling machine, and its surface was then ultrasonically cleaned in acetone for 10 min. To prepare the cladding material, equimolar NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW multicomponent powders were individually mixed using an attritor for 3 h. The mixed powders were then blended with an appropriate amount of an aqueous solution of 4 wt.% polyvinyl alcohol to form pieces of semisolid paste, each with a size of 120 × 6 × 2 mm. The semisolid paste was pre-placed on the surface of each AISI 1020 substrate, which was baked for 24 h at 150 °C to completely remove water and increase the adhesion between the cladding material and the substrate. After that, the cladding process was conducted. Table 1 indicates the welding parameters, which were suitable for the surface cladding of equimolar NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW multicomponent powders. After the cladding process, the various cladding layers were cut to suitable sizes for metallurgical
characterization
and
wear testing;
they were cut
by using wire
electrical-discharge machining (WEDM). An image analyzer was used to calculate the dilution rate of the cladding layer according to the cross-section of the cladding. The dimensions of the wear test specimen were controlled at 22 mm × 6 mm × (2 ± 0.01) mm, and an arc of radius 5 mm was fabricated by WEDM at the wear test end of the specimen to obtain the wear test geometry of the line-contact type. The wear test specimens were not subjected to 5
any heat treatment; hence the wear performance was assessed in the as-clad condition. The surface for the wear test was polished to control the surface roughness, making it be consistent in each specimen. The wear test was performed according to ASTM G99 standard [30], pin-on-disk mode with a rotating tribometer. The counterpart (disk) for the wear test was made of hardened AISI 52100 steel with a hardness of 61 HRc (≈ 720 Hv0.3) and a surface roughness controlled to approximately Ra = 0.06 µm. Each wear test under a particular condition was conducted at least four times to confirm the repeatability of the experiments. Table 2 presents the wear test conditions. After wear testing, the sample was examined by a stereo optical microscope used to measure the dimensions of the wear scars for calculating the wear loss of the wear test specimens. Nano-indentation was conducted to identify the micromechanical properties of the constituent phases in the cladding layer. Microhardness testing, x-ray diffraction (XRD), transmission electron microscopy (TEM), and electron probe microanalysis (EPMA) were used to determine the hardness distribution of the cladding layer, crystalline structure of the constituent phases in the cladding layer, microstructure of the cladding layer, and composition of the constituent phases, respectively. The XRD scanning angle, 2θ, was between 20° and 100°, and the scanning speed was 1°/min; the maximum acceleration voltage of TEM was 300 kV, and EPMA was 30 kV. Selected area diffraction (SAD) analysis was used to identify the crystal structure of the reinforcing phases and matrix in the cladding layer. 6
3. Results and discussion 3-1 Microstructural analyses of cladding layers In the multicomponent alloy-cladding process with the GTAW method, input energy per unit length less than the critical value resulted in the cladding layer failing to bond well on the substrate, and the cladding layer spalled. Nevertheless, the high input energy caused a high dilution rate of the cladding layer. Those phenomena might be attributable to the differences in the melting points of the equimolar components and low heat conductivity of the multicomponent alloys. The image analysis results indicated that the dilution rates of the NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW cladding layers were 64% and 51%, respectively. 3-1-1 Metallurgical characterization of the NiCrAlMoTiFeNbCo cladding layer Figure 1 displays the microstructure of the NiCrAlMoTiFeNbCo cladding layer at various positions. The reinforcements included some particulate and vein-shaped precipitates. In the cladding process, the cooling rates varied with the position, causing the microstructure in the cladding layer to vary. During the solidification process, the high cooling rate of the liquid metal resulted in a high degree of supercooling, and the critical nucleus radius was reduced. That caused the number of the nuclei to be increased in the region of the high cooling rate, and the grain size was smaller than that of the low cooling rate. Additionally, the thermal gradient and the solidification rate are the most important factors that determined the
7
grain shape and grain size [31]. Figure 1(a) shows the microstructure at the top of the cladding layer contained a larger number of vein-shaped reinforcements than at other positions; this was because the cooling rate there was higher. The reinforcements in the bottom region were sparser because of the dilution effect of the substrate. However, the interface between the cladding layer and the substrate exhibited a good metallurgical bonding, as shown in Fig. 1(c). The cooling rate in the middle region of the cladding layer was lower than that in all other regions, causing the microstructure there to be coarse. Figures 2 and 3 respectively display the results of the EPMA and XRD analyses of the NiCrAlMoTiFeNbCo cladding layer. According to these results, the matrix was α-Fe; the particulate reinforcements were composed of TiC, (Nb,Ti)C, and NbC, forming multiple-carbide particles with TiC-rich cores; and the vein-shaped reinforcements were composed of the eutectic phase Fe0.875Mo0.125. The EPMA results also indicated that Nb element was present in the Fe0.875Mo0.125 region. Moreover, a few particles of Al2O3 were detected in the cladding layer, and the related research indicated that the Al2O3 particles could contribute at the sites to heterogeneous nucleation of TiC [32], as shown in Fig. 2. A previous investigation demonstrated that the Ti/Nb atomic ratio dominated the formation of carbides when the melting pool temperature was below 2000 K. At temperatures below 2000 K, the Gibbs free energy of TiC (∆GTiC) was less than that of NbC [33]. Therefore, TiC particles formed first in the melting pool, reducing the Ti concentration in a microscale zone around the 8
TiC particles and changing the Ti/Nb ratio such that NbC was formed on the surfaces of the TiC particles. Thus, multiple-carbide particles (i.e., TiC particles wrapped in NbC) were formed. (Ti,Nb)C was also formed at the TiC/NbC interface. The EPMA and XRD analysis and SEM microstructural observation results confirmed this finding. Furthermore, the BEI image of the SEM in Fig. 1(a) revealed that the cores of the reinforcing particles were black, but the rest was grey, indirectly verifying that the reinforcing particle was a multiphase carbide. This result may be caused by the high cooling rate during solidification in the cladding process, which was different from the low cooling rate in the casting process [33, 34]. Figure 4 displays the TEM characterization results of the cladding layer. The reinforcing particles were distributed in the matrix, as depicted in Figs. 4(a) and (b). Figures 4(c) and (d) show the results of the selected area diffraction (SAD) analyses of the vein-shaped precipitates and the matrix, respectively, revealing that their crystal structures were BCC, but their lattice constants were 0.3046 nm for Fe0.875Mo0.125 and 0.2845 nm for α-Fe.
3-1-2 Metallurgical characterization of the NiCrAlMoTiFeNbW cladding layer Figure 5 shows the microstructure of the NiCrAlMoTiFeNbW cladding layer at various positions. The reinforcements included some particulates and petal-shaped precipitates. The microstructure at the top of the cladding layer contained more and smaller petal-shaped 9
precipitates compared to that at the other positions because the cooling rate was higher there. The dilution effect of the substrate caused the petal-shaped precipitates at the bottom region to be sparse and small in the NiCrAlMoTiFeNbW cladding layer. A few vein-shaped precipitates were also found in the bottom region of the NiCrAlMoTiFeNbW cladding layer. The interface between the bottom of the cladding layer and the substrate exhibited favorable metallurgical bonds. The cooling rate was the lowest in the middle of the cladding layer, causing the petal-shaped precipitates in this region to be coarse. Figures 6 and 7 show the EPMA and XRD results of the NiCrAlMoTiFeNbW cladding layer, respectively. According to these results, the matrix was α-Fe; the reinforcement comprised of TiC, (Nb,Ti)C, and NbC multiple-carbides; and the petal-shaped precipitates were composed of the Fe2W phase. A few Al2O3 particles were also detected in the cladding layer. The EPMA characterization showed that the Fe2W phase contained some Mo and Nb elements. The particulate reinforcements were composed of multiple carbides, and they had a TiC core and an NbC outer layer. The EPMA and XRD analyses and the SEM microstructural observations verified this finding. In the BEI image of the SEM in Fig. 5(a), the cores of the reinforcing particles were black, but the rest were gray, indicating that the compositions of the core and the outer layer differed. Figure 8 shows the TEM results of the NiCrAlMoTiFeNbW cladding layer. The reinforcing phases were distributed in the matrix, which is depicted in Fig. 8(a). Figures 8(b) 10
and (c) present the results of the SAD analyses of the reinforcement and the matrix, respectively. Figure 8(b) reveals that the crystalline structure of the reinforcement (Fe2W) was HCP with lattice constants of a = 0.5093 nm and c = 0.7470 nm. Figure 8(c) displays that the crystalline structure of the matrix was BCC with a lattice constant of a = 0.2751 nm. 3-2 Mechanical properties of the cladding layers The hardness (H) and Young’s modulus (E) of the cladding layers were characterized using a nanoindenter, and then the ratio of the H/E of each phase was calculated. Table 3 indicates that the hardness of the reinforcement Fe2W in the NiCrAlMoTiFeNbW cladding layer exceeded that of Fe0.875Mo0.125 in the NiCrAlMoTiFeNbCo cladding layer. Nevertheless, the hardness of the matrix of the NiCrAlMoTiFeNbW cladding layer was less than that of the matrix of the NiCrAlMoTiFeNbCo cladding layer. The EPMA analyses demonstrated that the matrix of the NiCrAlMoTiFeNbCo cladding layer contained Ni, Cr, Mo, and Co, while that of the NiCrAlMoTiFeNbW cladding layer contained Ni, Cr, Mo, and a little W. Furthermore, the lattice constant of the matrix of the NiCrAlMoTiFeNbCo cladding layer was greater than that of the matrix of the NiCrAlMoTiFeNbW cladding layer. These results implied that the solid solution strengthening effect of Co in the matrix was stronger than that of W in this cladding process. In contrast, the NiCrAlMoTiFeNbW cladding layer contained more of the reinforcing phases compared to the NiCrAlMoTiFeNbCo cladding layer because of the element W. This phenomenon was also related to the dilution rate of the NiCrAlMoTiFeNbW cladding layer, 11
which was less than that of the NiCrAlMoTiFeNbCo cladding layer. Therefore, the microhardness of the NiCrAlMoTiFeNbW cladding layer exceeded that of the NiCrAlMoTiFeNbCo cladding layer. Figure 9 presents the distribution of the microhardness from the cladding surface to the substrate, at a load of 300 g. 3-3 Tribological behaviors of the cladding layers
Figure 10 shows the wear loss of each specimen at different sliding speeds. The wear test results revealed that the wear resistance of all cladding specimens obviously improved, especially at a high sliding speed, compared to that of the AISI 1020 steel substrate. The wear loss of AISI 1020 steel at the high sliding speed was more than double that at low sliding speed because the former produced a high temperature in the contact region, and this high temperature led to the softening of the AISI 1020 steel along with plastic deformation accompanied by severe wear.
Figure 11 shows a typically worn surface of the AISI 1020 steel at a sliding speed of 1.643 m/s. The worn surface had the characteristics of plastic flow accompanying adhesive wear. Initially, the line contact mode caused high contact stress, higher than the yield stress of the low-carbon steel. Therefore, the real contact stress caused the high-stress asperities to experience plastic deformation, in which the accompanying shear effect led to adhesive wear behavior. After that, the apparent contacting area increased gradually with the wear, and the metal transfer and accumulated debris made scratches on the worn surface during sliding. 12
EDS characterization confirmed that the transfer film contained the elements from the counterpart and some oxides.
Figure 12 displays typical fluctuation of the friction coefficient during the wear test at high sliding speed for all the test specimens. The friction coefficient varied over a range of approximately 0.45–0.65. Moreover, the frictional behaviors of the wear test specimens at the low sliding speed were similar to that of the wear test specimens at the high sliding speed.
Figure 13 presents a typically worn surface of the NiCrAlMoTiFeNbCo-clad specimen, in which some patches of a “glaze” layer were formed. A previous study indicated that it is an oxide film [35]. The effect of precipitation strengthening of the reinforcements and solid solution strengthening reduced the plastic flow and adhesive behavior of the cladding layer that resulted in “glaze” oxide layer formation on the high-temperature rubbing surface and no severe wear on the sliding surface. However, the wear behavior during the sliding process caused the detachment of particulate reinforcements and made some small pits on the worn surface. The positions of the detached particulate reinforcements are clearly showing in the BEI image, Fig. 13(b), as black points. The detached particles contributed as the third bodies on the sliding surface, accompanying oxide film formation to reduce the friction coefficient [36]. Therefore, the friction coefficient of the NiCrAlMoTiFeNbCo-clad specimen was lower than that of the AISI 1020 steel specimen, varying over the range of approximately 0.4–0.6.
13
Additionally, the vein-shaped reinforcements improved the wear performance effectively because they can bond strongly on the sliding surface by mechanical locking due to complex geometric characteristics.
The NiCrAlMoTiFeNbW cladding specimen had the best wear performance of any of the wear test specimens, and the friction coefficient of that varied over a range of approximately 0.4–0.58. Figure 14 shows the worn surface of the NiCrAlMoTiFeNbW cladding specimen subjected to a low sliding speed. The worn surface displays petal-shaped precipitates that had a uniform distribution in the worn surface, bonding well with the matrix to resist the penetration of the counterpart and thus to improve the wear performance. Moreover, the thin “glaze” oxide film covered the worn surface completely. Figure 14(b) has a BEI image that indicates these petal-shaped precipitates distributed themselves on the worn surface more clearly, and the EDS semi-quantitative analysis (the diamond mark in Fig. 14 (b)) results indicated that worn surface contained oxygen, confirming there was a “glaze” oxide layer on the worn surface [37]. Mild abrasive wear with oxidation wear characteristics was observed, and the reinforcing phase Fe2W emerged on the worn surface. The Fe2W phase exhibited strong bonding with the matrix; it did not detach from the matrix during sliding, so the wear resistance of the cladding layer was improved effectively. Therefore, at a low sliding speed, the wear loss of the NiCrAlMoTiFeNbW-clad specimen was approximately 25% of that of the NiCrAlMoTiFeNbCo-clad specimen; however, at a high sliding speed, it was approximately 14
76%.
Figure 15 displays the worn surface of the NiCrAlMoTiFeNbW cladding specimen obtained after it was subjected to a high sliding speed. The shallow abrasive wear characteristics were observed, and some cracks were found in the petal-shaped reinforcing phase, which had the highest hardness than the other reinforcement, as shown in Table 3. At a high sliding speed, the temperature at the rubbing surfaces was high, leading to high thermal stress in the reinforcement, caused by the difference of thermal conductivity between matrix and reinforcement. Meanwhile, the traction force promoted the cracks to propagate normal to the sliding direction, resulting in thermal fatigue and detachment. As a result, the wear rate increased at the high sliding speed condition. In contrast, the wear loss of the NiCrAlMoTiFeNbCo cladding specimen at the high sliding speed was less than that at the low sliding speed because at the high sliding speed, the oxide film growth rate was high, which could reduce the adhesive behaviors. In all cladding specimens, the friction coefficient at the high sliding speed was less than that at the low sliding speed because the lower-friction oxide film formed quickly on the rubbing surface at the high temperature.
According to the experimental results, the NiCrAlMoTiFeNbW cladding layer exhibited excellent sliding wear resistance under the dry sliding wear test condition because Fe2W precipitate had good bonding with the matrix in the cladding layer, considerably improving
15
the wear resistance of the cladding layer. A previous investigation demonstrated that the ratio of hardness to Young’s modulus (H/E) significantly affects the wear resistance of a nanocomposite coating [38]. High H/E was associated with favorable wear performance. Table 3 shows the H/E values for the major phases in the clad specimens. The H/E ratios of both the matrix and the reinforcing phase Fe2W in the NiCrAlMoTiFeNbW cladding layer were higher than those of both the matrix and the reinforcing phase Fe0.875Mo0.125 in the NiCrAlMoTiFeNbCo cladding layer, respectively. For a fixed specimen hardness, a low Young’s modulus is associated with a high H/E ratio; therefore, at a given contact load, a lower E results in a higher elastic strain of the surface asperities, providing more contact positions and a larger real contact area. This effectively reduced the real contact stress and wear behavior. Accordingly, the NiCrAlMoTiFeNbW cladding layer had the best wear performance among the wear test specimens herein.
4. Conclusion The following conclusions are drawn from the results of the microstructural analysis and the wear testing: 1. The reinforcing particles formed in situ during surface cladding in both the NiCrAlMoTiFeNbCo and NiCrAlMoTiFeNbW cladding layers, comprised multiple layers of phases (Nb,Ti)C/NbC with a TiC-rich core. 2. The vein-shaped reinforcing phase Fe0.875Mo0.125 in the NiCrAlMoTiFeNbCo cladding layer and the petal-shaped reinforcing phase Fe2W in the NiCrAlMoTiFeNbW cladding layer were formed in situ during the cladding process, thereby increasing the hardness and improving the wear performance of the cladding layers. 16
3. The NiCrAlMoTiFeNbW cladding layer exhibited the best wear performance because this cladding layer had the highest H/E ratio and the Fe2W reinforcement phase was significantly harder than the Fe0.875Mo0.125 reinforcement phase. 4. The matrices of the different cladding layers had the same BCC crystal structure but different lattice constants because of solid solution strengthening containing different elements. 5. The multicomponent cladding process improved the tribological performance of the low-carbon steel. This process can extend the application of low-carbon steel to the moving components of a machine.
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20
Appendix: Wear volume calculation for the pin specimen
θ1 O
θ2 C
The shadow plane represents the worn surface of the pin specimen with a cylindrical end with a radius of curvature R Definition of symbols: L0: the width of the pin specimen R: radius of curvature of the cylindrical end θ1: semi-angle of the center on the small worn side θ2: semi-angle of the center on the large worn side α: semi-angle of the center at an arbitrary position of longitudinal cross section of the worn plane H1: width of the wear on the small worn side H2: width of the wear on the large worn side H: width of the wear of the longitudinal cross section of the worn plane at arbitrary position A: the area of the longitudinal cross section of the worn volume at an arbitrary position V: wear volume of the pin specimen Assume H2 ≧ H1, and θ2 ≦ 90∘ 21
Then,
，
,
If 1 = 2, 1 = 2 = , then the wear volume of the pin specimen can be calculated as follows. assume
，let
，
Then the term in the parentheses of the last term of Equation (A1) can be expressed as follow.
22
Use L'Hospital's Rule for eq. (A4)
∴ when
,
≈0
Note: the angle must be expressed using the unit of radian
23
Table 1 cladding parameter of each specimen Cladding materials
NiCrAlMoTiFeNbCo
NiCrAlMoTiFeNbW
Welding current (A)
110
Welding voltage (V)
17
Traveling speed(cm/min)
6.0
5.5
Flow rate of Ar (L/min)
10
Electrode
W-2% thorium
Polarity of electrode
DCEN (DCSP)
Diameter of W electrode (mm) Gap between electrode and workpiece (mm) Input energy per traveling length (kJ/cm)
1.6 2 18.7
20.4
Table 2. wear specimens and test conditions Dimension of the specimens (mm) Pin /Disk
Elastic modules (GPa)
22
6
2 with R5
186
NiCrAlMoTiFeNbW
186
AISI 52100 steel
210
0.29 / 0.3
(cladding layer / AISI 52100) Load (kg)
3.5
Maximum contact stress (MPa)
343
Sliding distance (m) Temperature Lubrication
11.5
NiCrAlMoTiFeNbCo
Poisson ratio
Sliding speed (m/s)
φ60
0.817
1.634 1470 25 Dry sliding
Table 3. Mechanical properties and the ratio of hardness to Young’s Modulus of each phase in the clad specimens NiCrAlMoTiFeNbCo clad layer Phases
Hardness (GPa) Young’s modulus(GPa)
H/E
Fe0.875Mo0.125 reinforcement
10.46
225
0.0465
Matrix
7.38
197
0.0375
NiCrAlMoTiFeNbW clad layer Fe2W reinforcement
16.34
246
0.0664
Matrix
7.20
186
0.0387
* Young’s modulus and hardness were inspected by nanoindenter * indentation load 2500 µN
Highlights 1. Equimolar multicomponent alloyed in-situ on steel to improve its wear performance. 2. Formation mechanism of the reinforcing phases in the clad layer was explored. 3. The correlation between the wear of the sample and the index of H/E was discussed. 4. Alloy element effect on the wear resistance of the clad layer was discussed.
Conflicts of Interest Statement
Manuscript title: Effect of Co and W on microstructure and wear behavior of NiCrAlMoTiFeNbX equimolar multicomponent clad layers Ref: WEA_2019_894
The authors whose names are listed immediately below certify that they have NO affiliations with or involvement in any organization or entity with any financial interest (such as honoraria; educational grants; participation in speakers’ bureaus; membership, employment, consultancies, stock ownership, or other equity interest; and expert testimony or patentlicensing arrangements), or non-financial interest (such as personal or professional relationships, affiliations, knowledge or beliefs) in the subject matter or materials discussed in this manuscript.
Author names: Yuan-Ching Lin Yu-Yu Liu