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Microstructure and wear behaviour of Ni-based surface coating on copper substrate
Wear 262 (2007) 868–875
Microstructure and wear behaviour of Ni-based surface coating on copper substrate Wen-ming Song a,b , Gui-rong Yang a,b , Jin-jun Lu c , Yuan Hao a,b , Ying Ma a,b a
State Key Laboratory of Gansu Advanced Non-ferrous Metal Materials, Lanzhou University of Technology, Lanzhou 730050, Gansu, China b Key Laboratory of Non-ferrous Metal Alloys, The Ministry of Education, Lanzhou University of Technology, Lanzhou 730050, China c State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, Gansu, China Received 4 January 2006; received in revised form 29 August 2006; accepted 31 August 2006 Available online 16 October 2006
Abstract The Ni-based surface coatings were prepared by a vacuum infiltration casting technique on copper substrate. The surface coatings were fabricated through copper melt penetrating into thin preforms whose thickness could change. By optimizing the processing parameters, compact surface coatings were achievable as confirmed through SEM observation. The surface coating was mainly composed of solid solution of Ni, solid solution of Cu and CrB. The macro-hardness of the coating was about HRC 58, and the micro-hardness of the coating shows a gradient distribution. The average micro-hardness of the coating was about HV450. Wear behaviour was investigated by using block-on-ring dry sliding linear contact at several loads (50 N–300 N) and two different sliding speeds (0.424 m/s and 0.848 m/s). Wear rate and friction coefficient were estimated using a method founded upon the PV factor theory. The surface oxidation predominated as the principle wear mechanism at low load. Meanwhile, adhesion and oxidation mechanism were observed when the coatings were tested at higher load more than 200 N. Friction coefficient decreased with increasing load and sliding speed. © 2006 Elsevier B.V. All rights reserved. Keywords: Copper alloy; Ni-based powder; Vacuum infiltration casting technique; Dry sliding wear; Wear rate
1. Introduction Surface condition of structural components has been a persistent problem in modern machinery. Some components stop functioning only due to the minor damage on their surface. By using coating technologies, it is possible to improve surface properties, such as wear, corrosion and oxidation resistances, and to take the advantage of a longer service life of coatings and the consequent reduction of the total cost. To this end, numerous surface modification techniques including thermal spraying, surface overlaying, chemical and physical deposition, high-energy beam cladding and infiltration technique have been widely used in industries [1–7]. Each coating technique has its own specified application area. A Thermally spraying layer is a mechanically bonded coating and the strength between coating and matrix is approximately 30–50 MPa at best. Surface overlaying produces a strong metallurgical bonding but at the same time induces high
E-mail address: [email protected] (G.-r. Yang). 0043-1648/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2006.08.025
welding residual stress and deformation that limit the shape and size of the parts. In addition, the smoothness of the overlaid surface is generally poor. Among these methods, the infiltration technique [6,7] has the advantages of simplicity and feasibility over others in changing the characteristic of copper surface. Infiltration of a liquid metal into a preform can be carried out by pressureless or pressure-assisted methods. The main advantages of pressure-assisted infiltration include the following: (a) faster process; (b) near-net shape processing; (c) reduced contamination; (d) reduced porosity in comparison with spray techniques [8–14]. Nickel-based alloys, used either on their own or combined with other reinforcement particles, have become popular because of both their outstanding wear and corrosion resistance at high temperatures and their relatively low cost [15–18]. Copper with high electrical and thermal conductivity is widely used in electronics and metallurgical industry; however, high strength and wear resistance are required in some applications. In order to improve wear resistance of copper, many methods have been developed as above mentioned. It is very difficult for
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fabricating surface wear resistant materials through infiltration technique on copper substrate because of its excellent thermal conductivity. In this study, a metallurgically bonded surface Ni-based coating on a copper substrate was prepared by vacuum infiltration casting technique. The microstructure and wear behaviour were investigated. 2. Experimental procedures 2.1. Fabrication of the surface coatings Ni-based alloying powder (Table 1) with particle size of 120–160 mesh and a spherical grain shape was used as the raw material of the surface coatings. Bronze ZQAl 9–4, with its composition listed in Table 2, was chosen as the matrix of the surface coatings and the substrate material on which the surface coatings formed. NJB (its main composition is boric acid). Coatings were fabricated by a vacuum infiltration casting technique. First, the Ni-based alloying powder and some binder were mixed together according to optimum proportion, paved on the inner surface of a casting mould or the outer surface of mould cores, where the surface strength and wear resistance of the components need improvement. Then the mould or mould cores was heated to form a preform layer at 150 ◦ C under the condition of atmosphere. The binder was fabricated by us and its main composition was boric acid. Second, the melt of bronze was poured into the mould at 1220 ◦ C in vacuum. The mould was to be preheated to 150 ± 10 ◦ C. By the sucking force of the vacuum, the melt was infiltrated into a preform, and surface coatings were obtained after solidification of the casting. The fabrication principle is shown in Fig. 1. Coatings
Table 1 Chemical composition of the Ni-based alloy powder (wt%) C B Si Cr Fe Mo Ni
0.6–0.8 2.5–3.5 4.0–5.0 16–18 ≤5.0 2.0–5.0 Bal.
Table 2 Chemical composition of tested aluminum bronze (wt%) Cu Al Fe Sb Si P As Sn Pb Mn Zn Total
Bal. 8.0–10.0 2.0–4.0 ≤0.05 ≤0.2 ≤0.1 ≤0.05 ≤0.2 ≤0.1 ≤0.5 ≤1.0 Bal.
Fig. 1. Schematic drawing for forming the Ni-based surface coatings.
with differential thickness can be obtained through preparing preforms with differential thickness. 2.2. Characterization and wear testing The specimens with a coating were cut, and the cross section was ground and polished using standard metallographic techniques for micro-structural observation. The morphology and thickness of the coatings were determined on the polished cross sections by optical and scanning electron microscopy HITACHI S-520 (S.E.M.). The composition was determined by an X-ray diffractometer (D/mac-rc) in the range of 30–100◦ at the speed of 4◦ /min. The macro-hardness of the coatings was measured using HRS-150 Rockwell hardness tester under 1471 N load, and its micro-hardness was measured using HVS-1000 micro-hardness tester at a constant load of 9.8 N and dwelling time of 20 s. The wear behaviour of coating was evaluated without lubrication at room temperature using a block-on-ring tester. The wear tests were carried out on MM-2000 wear test machine. The block was the specimen with surface coatings, and its dimension was 6 mm × 7 mm × 30 mm. The surface of coatings for wear tests was ground, ultrasonically cleaned in water, and then rinsed in ethanol prior to drying in air. It was rectified (average roughness Ra = 0.74 m) and placed against the counterbody, a hardened and tempered GCr15 steel ring, with an outer diameter of 40 mm and thickness of 10 mm, the average surface macrohardness of the ring is HRC 64. This setup was used to analyze the effect of load and sliding speed on the wear resistance of the coatings. A number of wear tests were conducted at different loads (50 N, 100 N, 200 N, 300 N) at two constant sliding speeds (0.424 m/s and 0.848 m/s), respectively. Three tests were run for each of the specified test conditions. The sliding distance was 1500 m. Coatings and worn surfaces were analyzed by optical microscopy and scanning electron microscopy with energy dispersive spectroscopy (EDS). Friction coefficient was determined using a method founded upon the PV factor theory [19–21]. Data was collected through a computer. The width of the wear track was also measured at the end of each test so as to calculate the total wear volume. The following formula was
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used to estimate the wear rate. S = 4.67 × b × (R − R × R − b × b/4), W=
L = v × t,
S L
where b is the width of wear track (mm), and R is radius of the counterbody (mm), and L is the distance of sliding (m), v is the line velocity (m/s), t is the time of sliding (s), W is wear rate (mm3 /m). The surface temperature of specimen was measured using a thermocouple with a precision of 1 ◦ C. The measurement area was located at the dot of the specimen that was 2.0 mm below the wear contact interface. 3. Results and discussion 3.1. Microstructure of the surface coatings The cross-sectional views of the surface coatings are shown in Figs. 2 and 3. The micrographs reveal continuous and sound interface between the substrate and the coating without crack present. No spalling or separation is observed even at high magnification (1000×). Fig. 2 shows that coating consists of three parts: a surface sintering layer (SSL), a metallurgical fusing layer (MFL) and a diffusion solid solution layer (DSSL). Fig. 3 shows the morphology of SSL, which indicate that the microstructure was compact. The surface coating formation was a kind of metallurgical fusion because the melting point of Ni-based alloying powder is about 1050 ◦ C lower than the pouring temperature of the copper alloy. After pouring the melt, the metal infiltrated into the pores and permeated through the whole preform to the outer layer of the preform close to the mould surface. In the outer-
Fig. 2. Cross-sectional S.E.M. view of the coating.
Fig. 3. S.E.M. view of the outer layer.
most layer, the Ni-based alloying grains partially melt, and the copper melt could only sinter with the alloying powder at the original site. The composition of the outer surface sintering layer was mainly that of Ni-based alloy with a small amount of copper. The temperature of the following copper melt for forming MFL was higher than that of the copper melt for forming the SSL. Therefore, metallurgical fusion happened between the copper alloy and Ni-based powder to cause the formation of the metallurgical fusion layer. The Ni-based powder for forming the innermost layer of the whole surface coating obtained the most energy compared with that for forming the outer two layers. The Ni-based alloying powder could melt, decompose and diffuse only to form many active atoms that may infiltrate directly into the substrate to form solid solution, which is the diffusion solid solution layer. The distributions of Ni, Cr, and Cu across the substrate and the surface coating were continuous (shown in Fig. 4). The distribution of Ni and Cr gradually decreases from the outer layer to the inner layer. On the other hand, the distribution of Cu increases gradually from the outer layer to the inner layer. The total thickness of DSSL and MFL was larger than that obtained through plating, spraying and vapor deposition because of faster diffusion in liquid and longer diffusion distance [22–24]. The diffusion temperature of other methods, such as heat treatment (generally at 800 ◦ C for about 24 h), was lower than that of the vacuum infiltration casting technique. The structure of the surface material is promising from the point view of microstructure of structural materials because it has a transition layer (DSSL). Combining the observations from Figs. 4 and 5, the following conclusions could be drawn. The outer layer (SSL) was mainly
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Fig. 6. The variation of micro-hardness.
Fig. 4. Line scan of element distribution across the whole coating.
composed of CrB and a solid solution of Ni. The metallurgical fusing layer (MFL) was mainly composed of CrB, a solid solution of Cu and solid solution of Ni. The inner layer (DSSL) is mainly composed of a solid solution of Cu and some solid solutions of Ni. 3.2. Hardness of surface coatings The hardness of the copper alloy is around HRB 130, while, the hardness of the coating is about HRC58. The hardness of surface coating had been improved significantly. Fig. 6 shows the variation of micro-hardness across the whole surface coating. The outer layer near the casting mould contains unavoidably shrinkage porosity because of the chilling action that leads to
Fig. 5. XRD of the surface coating.
imperfect fusion. The compactness of the outer layer was not as good as that of the sub-surface layer, and the hardness of the outer layer was not the highest. The hardness reached the highest value in the compact sub-surface layer because the quality of fusion was better there than that in the outer layer, moreover, it contained more hard phases. The active atom decomposing from the melt of the Ni-based alloying powder located at the innermost part of the preform diffused into the Cu substrate to form the diffusion area with considerable thickness. The hard phase could not diffuse to this layer so its content drops in this region. The solid solution of Cu is the main composition of this layer. That is the main factor for apparent decreasing in microhardness in this layer. 3.3. Wear test results Fig. 7 shows the wear rate of the coating as a function of load at two different sliding speeds. Wear rate increased with the increasing sliding speed and load applied on the specimen. The tests with varied sliding speed and constant load (50 N) indicated a minor influence of sliding speed on wear rate. The
Fig. 7. Wear rate of the coating as a function of load at two different sliding speeds.
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wear rate of coating at high sliding speed was twice as large as that at low sliding speed whilst the load remained constant ranged from 100 N to 200 N. The wear rate of coating at a 200 N load was also twice as large as that at 100 N whilst the sliding speed remained constant. The difference of wear rate between high sliding speed and low sliding speed became large with the increasing load, and it nearly kept constant when the load was beyond 200 N. The worn surface of the Ni-based surface coating is smooth, as shown in Fig. 8(a) and (b) after tested at a high sliding speed and low load. There is a layer of loose oxide on the worn surface. Fig. 8(c) is an EDS analysis. It illustrates that the content of element Fe and Cr was much more than that of element Ni, while the
Fig. 8. (a) S.E.M. photographs of the worn surface of a coating at the 100 N and 0.848 m/s; (b) higher magnification of (a) to show the clear surface appearance; (c) EDS analysis of the worn surface.
element Fe content of the coating was less than 5%. This observation suggests that some Fe and Cr be transferred from the GCr 15 steel ring to the surface of the specimen, and there was oxidation wear evidence. The oxide layer on the surface of the specimen seemed to form faster than it could wear up under the condition of low load and low sliding speed, thereby, leaving oxide layer on the worn surface of the specimen. This could decrease the wear rate of the coatings. On the contrary, the oxide layer on the surface of the specimen seemed to wear up faster than it could form when the sliding speed and load were high, therefore, leaving new material appearance to the interface between the specimen and the counterbody which could then oxidize and wear away. The wear course alternated between forming and removal of the oxide layer, which resulted in increased abrasion. In a word, the oxidation wear mechanism was the main wear mechanism under the low load condition when the load is lower than 200 N. Fig. 9 shows the worn surface of the specimen under the condition of high sliding speed and high load. The surface oxide layer of the specimen was destroyed indicated from Fig. 9(a)–(c). It was significant that more plastically deformed platelets containing Ni, Cr, Fe and Si (Fig. 10) and much more wear rust were found adhere to the wear track of the conterbody as the load was increased. This indicated that there was some adhesive wear combined with the oxidation process observed for the load higher than 200 N. The friction coefficient was also measured in each test and its value was calculated. Fig. 11 shows how the coefficient changed as the load was increased. Similar results were obtained for two different sliding speeds (low and high sliding speeds). The amplitude of friction coefficient became smaller when the load was increased. The course of friction coefficient becoming steady was shortened under the high sliding speed condition. Lower loads might also not as effective in removal of wear debris from the sliding interface, which probably increased the frictional resistance. This indicated that the thickness of wear debris was not uniform. However, the debris could be removed in time on high load, and the new coating materials surface appeared. The rust on the surface of the GCr 15 conterbody seems to have worked as a kind of lubricant. In a word, the lowering of friction coefficient with increasing load can be explained by the smoothening of the coating surface and the rust on the surface of the GCr 15 conterbody as a kind of lubricant. In the course of the wear test, surface oxidation and deformation of the counterbody accompanied friction heat generated simultaneously, which resulted in the hardness decrease of the GCr 15 counterbody. Fig. 12 indicates the surface temperature variation of the sample with the increasing PV value. The surface temperature increased in a linear relationship with the increasing PV value. The increasing surface temperature accelerated the surface oxidation and resulted in the tempering of the steel ring. The hardness of the GCr 15 counterbody was HRC 64 before wear test. It decreased to HRC 58. This result agrees with the PV factor theory described by Halling [19] and Cadenas et al. [20], who derived a similar expression based upon the hypothesis that wear rate is proportional to the power dissipated at the sliding interface.
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Fig. 9. (a) S.E.M. photographs of the worn surface of a coating tested under 300 N and 0.848 m/s; (b) higher magnification of area A in (a) to show the clear surface appearance; (c) higher magnification of area B in (a) to show the clear surface appearance; (d) EDS analysis of the worn surface.
The wear tests of the substrate material (copper alloy) were conducted at a sliding speed 0.848 m/s under load 200 N and 300 N, respectively. The wear rate was 1.12 × 10−2 mm3 /m and 2.85 × 10−2 mm3 /m, which was much higher than that of the
surface coatings. The temperature of the worn surface for substrate increased to 427 K and 489 K very fast, respectively. The friction heat resulted in softening of the substrate, which caused the strength of the substrate to decrease. There were more plas-
Fig. 10. (a) S.E.M. photograph of the worn surface of the counterbody at the load 300 N and sliding speed 0.848 m/s; (b) EDS analysis of the worn surface of the counterbody.
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Fig. 11. Variation in friction coefficient: (a) as a function of load at a constant sliding speed of 0.424 m/s and (b) as a function of load at a constant sliding speed of 0.848 m/s.
conditions. The fluctuating amplitude was decreased with the increasing load and the increasing sliding speed. Finally, the wear rate of the Ni-based surface coatings and the substrate differ by two orders of magnitude. Acknowledgements
Fig. 12. Variation in temperature measured under the coating in contact with a hardened, tempered GCr15 ring, as a function of PV product.
tically deformations at the surface of the substrate contacting with the counterbody. 4. Conclusion A vacuum infiltration casting technology can be used for the fabrication of surface coatings through the correct choice of parameters. Ni-based powder has been successfully used as the surface alloying powder to obtain good Ni-based surface coatings on a copper substrate. The surface coating has clearly excellent interfacial boundaries due to long-range diffusion of elements across the interface. The coating includes three sublayers: surface sintering layer, metallurgical fusing layer and diffusion solid solution layer. The outer layer was mainly composed of CrB and a solid solution of Ni. The metallurgical fusing layer was mainly composed of CrB, a solid solution of Cu and a solid solution of Ni. The inner layer was mainly composed of a solid solution of Cu. The macro-hardness of the coating reaches about HRC 58, and highest micro-hardness is at the sub-surface of the coating. A minor influence of sliding speed on wear rate was observed in the dry sliding block-on-ring wear tests at low load. Wear rate increased with the increasing of sliding speed and load. This process of wear by oxidation was maintained under all test conditions and was accompanied by adhesive mechanism at the high load. The increasing of surface temperature of the specimen was linear with the increasing PV value during the wear course. The friction coefficient could not keep constant under all test
The authors would like to acknowledge the financial support for this research provided by the Chun-Hui Plan of the Education Department of China project (Z2004-1-62013), and the Young Teacher Startup Foundation project at Lanzhou University of Technology. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
[18]
J. Choi, J. Mazumder, J. Mater. Sci. 29 (1994) 4460. W. Cerri, R. Maryinella, G.P. Mor, Surf. Coat. Technol. 49 (1991) 40. A. Ramalho, A. Merstallinger, A. Cavaleiro, Wear 260 (2006) 361. A. Ramalho, M.T. Vieira, Surf. Coat. Technol. 197 (2005) 358. Q. Ming, L.C. Lim, Z.D. Chen, Surf. Coat. Technol. 106 (1998) 174. Yang Gui-rong, Hao Yuan, Song Wen-ming, Ma Ying, Mater. Sci. Eng. A 399 (2005) 206. J. Lapin, D. Tiberghien, F. Delannay, Intermetallics 8 (2000) 1429– 1438. D.C. Dunand, Mater. Manuf. Processes 10 (1995) 373. H. Nakae, H. Fujii, K. Nakajima, A. Goto, Mater. Sci. Eng. A 298 (2002) 210. S. San Marchi, A. Mortensen, Composites Part A: Appl. Sci. Manuf. 34 (2003) 907. Y. Chen, D.D.L. Chung, J. Mater. Sci. 31 (1996) 2117. D.L. Joslin, D.S. Easton, C.T. Liu, S.A. David, Mater. Sci. Eng. A 192–193 (1995) 544. D.L. Joslin, D.S. Easton, C.T. Liu, S.S. Babu, S.A. David, Intermetallics 3 (1995) 467. L.Z. Zhuang, L. Buekenhout, J. Duszczyk, Script Metal Mater. 30 (1994) 909. L.C. Qian Ming, Z.D. Lim, Chen Laser cladding of nickel-based hardfacing alloys, Surf. Coat. Technol. 106 (1998) 174–182. A. Conde, F. Zubiri, J. Damborenea, Cladding of Ni-Cr-B-Si coatings with a high power diode laser, Mater. Sci. Eng. A 334 (1–2) (2002) 233–238. J. Rodr’iguez, A. Mart’in, R. Fern’andez, J.E. Fern’andez, An experimental study of the wear performance of NiCrBSi thermal spray coatings, Wear 255 (2003) 950–955. P. Wu, X.L. Chen, E.Y. Jiang, Influence of WC particle behavior on the wear resistance properties of Ni-WC composite coatings, Wear 257 (1–2) (2004) 142–147.
W.-m. Song et al. / Wear 262 (2007) 868–875 [19] J. Halling, Principle of Tribology, McMillan, London, 1975. [20] M. Cadenas, R. Vijande, H.J. Montes, J.M. Sierra, Wear behaviour of laser cladded and plasma sprayed WC-Co coatings, Wear 212 (1997) 244– 253. [21] E. Fernandez, M. Cadenas, R. Gonzalez, Wear 259 (2005) 870–875.
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[22] D.C. Dunand, Mater. Manuf. Processes 10 (1995) 373. [23] H. Nakae, H. Fujii, K. Nakajima, A. Goto, Mater. Sci. Eng. A 223 (1997) 21. [24] S. San Marchi, A. Mortensen, Metall. Mater. Trans. A 29 (1998) 2819.
Microstructure and wear behaviour of Ni-based surface coating on copper substrate Wen-ming Song a,b , Gui-rong Yang a,b , Jin-jun Lu c , Yuan Hao a,b , Ying Ma a,b a
State Key Laboratory of Gansu Advanced Non-ferrous Metal Materials, Lanzhou University of Technology, Lanzhou 730050, Gansu, China b Key Laboratory of Non-ferrous Metal Alloys, The Ministry of Education, Lanzhou University of Technology, Lanzhou 730050, China c State Key Laboratory of Solid Lubrication, Lanzhou Institute of Chemical Physics, Chinese Academy of Sciences, Lanzhou 730000, Gansu, China Received 4 January 2006; received in revised form 29 August 2006; accepted 31 August 2006 Available online 16 October 2006
Abstract The Ni-based surface coatings were prepared by a vacuum infiltration casting technique on copper substrate. The surface coatings were fabricated through copper melt penetrating into thin preforms whose thickness could change. By optimizing the processing parameters, compact surface coatings were achievable as confirmed through SEM observation. The surface coating was mainly composed of solid solution of Ni, solid solution of Cu and CrB. The macro-hardness of the coating was about HRC 58, and the micro-hardness of the coating shows a gradient distribution. The average micro-hardness of the coating was about HV450. Wear behaviour was investigated by using block-on-ring dry sliding linear contact at several loads (50 N–300 N) and two different sliding speeds (0.424 m/s and 0.848 m/s). Wear rate and friction coefficient were estimated using a method founded upon the PV factor theory. The surface oxidation predominated as the principle wear mechanism at low load. Meanwhile, adhesion and oxidation mechanism were observed when the coatings were tested at higher load more than 200 N. Friction coefficient decreased with increasing load and sliding speed. © 2006 Elsevier B.V. All rights reserved. Keywords: Copper alloy; Ni-based powder; Vacuum infiltration casting technique; Dry sliding wear; Wear rate
1. Introduction Surface condition of structural components has been a persistent problem in modern machinery. Some components stop functioning only due to the minor damage on their surface. By using coating technologies, it is possible to improve surface properties, such as wear, corrosion and oxidation resistances, and to take the advantage of a longer service life of coatings and the consequent reduction of the total cost. To this end, numerous surface modification techniques including thermal spraying, surface overlaying, chemical and physical deposition, high-energy beam cladding and infiltration technique have been widely used in industries [1–7]. Each coating technique has its own specified application area. A Thermally spraying layer is a mechanically bonded coating and the strength between coating and matrix is approximately 30–50 MPa at best. Surface overlaying produces a strong metallurgical bonding but at the same time induces high
E-mail address: [email protected] (G.-r. Yang). 0043-1648/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.wear.2006.08.025
welding residual stress and deformation that limit the shape and size of the parts. In addition, the smoothness of the overlaid surface is generally poor. Among these methods, the infiltration technique [6,7] has the advantages of simplicity and feasibility over others in changing the characteristic of copper surface. Infiltration of a liquid metal into a preform can be carried out by pressureless or pressure-assisted methods. The main advantages of pressure-assisted infiltration include the following: (a) faster process; (b) near-net shape processing; (c) reduced contamination; (d) reduced porosity in comparison with spray techniques [8–14]. Nickel-based alloys, used either on their own or combined with other reinforcement particles, have become popular because of both their outstanding wear and corrosion resistance at high temperatures and their relatively low cost [15–18]. Copper with high electrical and thermal conductivity is widely used in electronics and metallurgical industry; however, high strength and wear resistance are required in some applications. In order to improve wear resistance of copper, many methods have been developed as above mentioned. It is very difficult for
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fabricating surface wear resistant materials through infiltration technique on copper substrate because of its excellent thermal conductivity. In this study, a metallurgically bonded surface Ni-based coating on a copper substrate was prepared by vacuum infiltration casting technique. The microstructure and wear behaviour were investigated. 2. Experimental procedures 2.1. Fabrication of the surface coatings Ni-based alloying powder (Table 1) with particle size of 120–160 mesh and a spherical grain shape was used as the raw material of the surface coatings. Bronze ZQAl 9–4, with its composition listed in Table 2, was chosen as the matrix of the surface coatings and the substrate material on which the surface coatings formed. NJB (its main composition is boric acid). Coatings were fabricated by a vacuum infiltration casting technique. First, the Ni-based alloying powder and some binder were mixed together according to optimum proportion, paved on the inner surface of a casting mould or the outer surface of mould cores, where the surface strength and wear resistance of the components need improvement. Then the mould or mould cores was heated to form a preform layer at 150 ◦ C under the condition of atmosphere. The binder was fabricated by us and its main composition was boric acid. Second, the melt of bronze was poured into the mould at 1220 ◦ C in vacuum. The mould was to be preheated to 150 ± 10 ◦ C. By the sucking force of the vacuum, the melt was infiltrated into a preform, and surface coatings were obtained after solidification of the casting. The fabrication principle is shown in Fig. 1. Coatings
Table 1 Chemical composition of the Ni-based alloy powder (wt%) C B Si Cr Fe Mo Ni
0.6–0.8 2.5–3.5 4.0–5.0 16–18 ≤5.0 2.0–5.0 Bal.
Table 2 Chemical composition of tested aluminum bronze (wt%) Cu Al Fe Sb Si P As Sn Pb Mn Zn Total
Bal. 8.0–10.0 2.0–4.0 ≤0.05 ≤0.2 ≤0.1 ≤0.05 ≤0.2 ≤0.1 ≤0.5 ≤1.0 Bal.
Fig. 1. Schematic drawing for forming the Ni-based surface coatings.
with differential thickness can be obtained through preparing preforms with differential thickness. 2.2. Characterization and wear testing The specimens with a coating were cut, and the cross section was ground and polished using standard metallographic techniques for micro-structural observation. The morphology and thickness of the coatings were determined on the polished cross sections by optical and scanning electron microscopy HITACHI S-520 (S.E.M.). The composition was determined by an X-ray diffractometer (D/mac-rc) in the range of 30–100◦ at the speed of 4◦ /min. The macro-hardness of the coatings was measured using HRS-150 Rockwell hardness tester under 1471 N load, and its micro-hardness was measured using HVS-1000 micro-hardness tester at a constant load of 9.8 N and dwelling time of 20 s. The wear behaviour of coating was evaluated without lubrication at room temperature using a block-on-ring tester. The wear tests were carried out on MM-2000 wear test machine. The block was the specimen with surface coatings, and its dimension was 6 mm × 7 mm × 30 mm. The surface of coatings for wear tests was ground, ultrasonically cleaned in water, and then rinsed in ethanol prior to drying in air. It was rectified (average roughness Ra = 0.74 m) and placed against the counterbody, a hardened and tempered GCr15 steel ring, with an outer diameter of 40 mm and thickness of 10 mm, the average surface macrohardness of the ring is HRC 64. This setup was used to analyze the effect of load and sliding speed on the wear resistance of the coatings. A number of wear tests were conducted at different loads (50 N, 100 N, 200 N, 300 N) at two constant sliding speeds (0.424 m/s and 0.848 m/s), respectively. Three tests were run for each of the specified test conditions. The sliding distance was 1500 m. Coatings and worn surfaces were analyzed by optical microscopy and scanning electron microscopy with energy dispersive spectroscopy (EDS). Friction coefficient was determined using a method founded upon the PV factor theory [19–21]. Data was collected through a computer. The width of the wear track was also measured at the end of each test so as to calculate the total wear volume. The following formula was
870
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used to estimate the wear rate. S = 4.67 × b × (R − R × R − b × b/4), W=
L = v × t,
S L
where b is the width of wear track (mm), and R is radius of the counterbody (mm), and L is the distance of sliding (m), v is the line velocity (m/s), t is the time of sliding (s), W is wear rate (mm3 /m). The surface temperature of specimen was measured using a thermocouple with a precision of 1 ◦ C. The measurement area was located at the dot of the specimen that was 2.0 mm below the wear contact interface. 3. Results and discussion 3.1. Microstructure of the surface coatings The cross-sectional views of the surface coatings are shown in Figs. 2 and 3. The micrographs reveal continuous and sound interface between the substrate and the coating without crack present. No spalling or separation is observed even at high magnification (1000×). Fig. 2 shows that coating consists of three parts: a surface sintering layer (SSL), a metallurgical fusing layer (MFL) and a diffusion solid solution layer (DSSL). Fig. 3 shows the morphology of SSL, which indicate that the microstructure was compact. The surface coating formation was a kind of metallurgical fusion because the melting point of Ni-based alloying powder is about 1050 ◦ C lower than the pouring temperature of the copper alloy. After pouring the melt, the metal infiltrated into the pores and permeated through the whole preform to the outer layer of the preform close to the mould surface. In the outer-
Fig. 2. Cross-sectional S.E.M. view of the coating.
Fig. 3. S.E.M. view of the outer layer.
most layer, the Ni-based alloying grains partially melt, and the copper melt could only sinter with the alloying powder at the original site. The composition of the outer surface sintering layer was mainly that of Ni-based alloy with a small amount of copper. The temperature of the following copper melt for forming MFL was higher than that of the copper melt for forming the SSL. Therefore, metallurgical fusion happened between the copper alloy and Ni-based powder to cause the formation of the metallurgical fusion layer. The Ni-based powder for forming the innermost layer of the whole surface coating obtained the most energy compared with that for forming the outer two layers. The Ni-based alloying powder could melt, decompose and diffuse only to form many active atoms that may infiltrate directly into the substrate to form solid solution, which is the diffusion solid solution layer. The distributions of Ni, Cr, and Cu across the substrate and the surface coating were continuous (shown in Fig. 4). The distribution of Ni and Cr gradually decreases from the outer layer to the inner layer. On the other hand, the distribution of Cu increases gradually from the outer layer to the inner layer. The total thickness of DSSL and MFL was larger than that obtained through plating, spraying and vapor deposition because of faster diffusion in liquid and longer diffusion distance [22–24]. The diffusion temperature of other methods, such as heat treatment (generally at 800 ◦ C for about 24 h), was lower than that of the vacuum infiltration casting technique. The structure of the surface material is promising from the point view of microstructure of structural materials because it has a transition layer (DSSL). Combining the observations from Figs. 4 and 5, the following conclusions could be drawn. The outer layer (SSL) was mainly
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Fig. 6. The variation of micro-hardness.
Fig. 4. Line scan of element distribution across the whole coating.
composed of CrB and a solid solution of Ni. The metallurgical fusing layer (MFL) was mainly composed of CrB, a solid solution of Cu and solid solution of Ni. The inner layer (DSSL) is mainly composed of a solid solution of Cu and some solid solutions of Ni. 3.2. Hardness of surface coatings The hardness of the copper alloy is around HRB 130, while, the hardness of the coating is about HRC58. The hardness of surface coating had been improved significantly. Fig. 6 shows the variation of micro-hardness across the whole surface coating. The outer layer near the casting mould contains unavoidably shrinkage porosity because of the chilling action that leads to
Fig. 5. XRD of the surface coating.
imperfect fusion. The compactness of the outer layer was not as good as that of the sub-surface layer, and the hardness of the outer layer was not the highest. The hardness reached the highest value in the compact sub-surface layer because the quality of fusion was better there than that in the outer layer, moreover, it contained more hard phases. The active atom decomposing from the melt of the Ni-based alloying powder located at the innermost part of the preform diffused into the Cu substrate to form the diffusion area with considerable thickness. The hard phase could not diffuse to this layer so its content drops in this region. The solid solution of Cu is the main composition of this layer. That is the main factor for apparent decreasing in microhardness in this layer. 3.3. Wear test results Fig. 7 shows the wear rate of the coating as a function of load at two different sliding speeds. Wear rate increased with the increasing sliding speed and load applied on the specimen. The tests with varied sliding speed and constant load (50 N) indicated a minor influence of sliding speed on wear rate. The
Fig. 7. Wear rate of the coating as a function of load at two different sliding speeds.
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wear rate of coating at high sliding speed was twice as large as that at low sliding speed whilst the load remained constant ranged from 100 N to 200 N. The wear rate of coating at a 200 N load was also twice as large as that at 100 N whilst the sliding speed remained constant. The difference of wear rate between high sliding speed and low sliding speed became large with the increasing load, and it nearly kept constant when the load was beyond 200 N. The worn surface of the Ni-based surface coating is smooth, as shown in Fig. 8(a) and (b) after tested at a high sliding speed and low load. There is a layer of loose oxide on the worn surface. Fig. 8(c) is an EDS analysis. It illustrates that the content of element Fe and Cr was much more than that of element Ni, while the
Fig. 8. (a) S.E.M. photographs of the worn surface of a coating at the 100 N and 0.848 m/s; (b) higher magnification of (a) to show the clear surface appearance; (c) EDS analysis of the worn surface.
element Fe content of the coating was less than 5%. This observation suggests that some Fe and Cr be transferred from the GCr 15 steel ring to the surface of the specimen, and there was oxidation wear evidence. The oxide layer on the surface of the specimen seemed to form faster than it could wear up under the condition of low load and low sliding speed, thereby, leaving oxide layer on the worn surface of the specimen. This could decrease the wear rate of the coatings. On the contrary, the oxide layer on the surface of the specimen seemed to wear up faster than it could form when the sliding speed and load were high, therefore, leaving new material appearance to the interface between the specimen and the counterbody which could then oxidize and wear away. The wear course alternated between forming and removal of the oxide layer, which resulted in increased abrasion. In a word, the oxidation wear mechanism was the main wear mechanism under the low load condition when the load is lower than 200 N. Fig. 9 shows the worn surface of the specimen under the condition of high sliding speed and high load. The surface oxide layer of the specimen was destroyed indicated from Fig. 9(a)–(c). It was significant that more plastically deformed platelets containing Ni, Cr, Fe and Si (Fig. 10) and much more wear rust were found adhere to the wear track of the conterbody as the load was increased. This indicated that there was some adhesive wear combined with the oxidation process observed for the load higher than 200 N. The friction coefficient was also measured in each test and its value was calculated. Fig. 11 shows how the coefficient changed as the load was increased. Similar results were obtained for two different sliding speeds (low and high sliding speeds). The amplitude of friction coefficient became smaller when the load was increased. The course of friction coefficient becoming steady was shortened under the high sliding speed condition. Lower loads might also not as effective in removal of wear debris from the sliding interface, which probably increased the frictional resistance. This indicated that the thickness of wear debris was not uniform. However, the debris could be removed in time on high load, and the new coating materials surface appeared. The rust on the surface of the GCr 15 conterbody seems to have worked as a kind of lubricant. In a word, the lowering of friction coefficient with increasing load can be explained by the smoothening of the coating surface and the rust on the surface of the GCr 15 conterbody as a kind of lubricant. In the course of the wear test, surface oxidation and deformation of the counterbody accompanied friction heat generated simultaneously, which resulted in the hardness decrease of the GCr 15 counterbody. Fig. 12 indicates the surface temperature variation of the sample with the increasing PV value. The surface temperature increased in a linear relationship with the increasing PV value. The increasing surface temperature accelerated the surface oxidation and resulted in the tempering of the steel ring. The hardness of the GCr 15 counterbody was HRC 64 before wear test. It decreased to HRC 58. This result agrees with the PV factor theory described by Halling [19] and Cadenas et al. [20], who derived a similar expression based upon the hypothesis that wear rate is proportional to the power dissipated at the sliding interface.
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Fig. 9. (a) S.E.M. photographs of the worn surface of a coating tested under 300 N and 0.848 m/s; (b) higher magnification of area A in (a) to show the clear surface appearance; (c) higher magnification of area B in (a) to show the clear surface appearance; (d) EDS analysis of the worn surface.
The wear tests of the substrate material (copper alloy) were conducted at a sliding speed 0.848 m/s under load 200 N and 300 N, respectively. The wear rate was 1.12 × 10−2 mm3 /m and 2.85 × 10−2 mm3 /m, which was much higher than that of the
surface coatings. The temperature of the worn surface for substrate increased to 427 K and 489 K very fast, respectively. The friction heat resulted in softening of the substrate, which caused the strength of the substrate to decrease. There were more plas-
Fig. 10. (a) S.E.M. photograph of the worn surface of the counterbody at the load 300 N and sliding speed 0.848 m/s; (b) EDS analysis of the worn surface of the counterbody.
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Fig. 11. Variation in friction coefficient: (a) as a function of load at a constant sliding speed of 0.424 m/s and (b) as a function of load at a constant sliding speed of 0.848 m/s.
conditions. The fluctuating amplitude was decreased with the increasing load and the increasing sliding speed. Finally, the wear rate of the Ni-based surface coatings and the substrate differ by two orders of magnitude. Acknowledgements
Fig. 12. Variation in temperature measured under the coating in contact with a hardened, tempered GCr15 ring, as a function of PV product.
tically deformations at the surface of the substrate contacting with the counterbody. 4. Conclusion A vacuum infiltration casting technology can be used for the fabrication of surface coatings through the correct choice of parameters. Ni-based powder has been successfully used as the surface alloying powder to obtain good Ni-based surface coatings on a copper substrate. The surface coating has clearly excellent interfacial boundaries due to long-range diffusion of elements across the interface. The coating includes three sublayers: surface sintering layer, metallurgical fusing layer and diffusion solid solution layer. The outer layer was mainly composed of CrB and a solid solution of Ni. The metallurgical fusing layer was mainly composed of CrB, a solid solution of Cu and a solid solution of Ni. The inner layer was mainly composed of a solid solution of Cu. The macro-hardness of the coating reaches about HRC 58, and highest micro-hardness is at the sub-surface of the coating. A minor influence of sliding speed on wear rate was observed in the dry sliding block-on-ring wear tests at low load. Wear rate increased with the increasing of sliding speed and load. This process of wear by oxidation was maintained under all test conditions and was accompanied by adhesive mechanism at the high load. The increasing of surface temperature of the specimen was linear with the increasing PV value during the wear course. The friction coefficient could not keep constant under all test
The authors would like to acknowledge the financial support for this research provided by the Chun-Hui Plan of the Education Department of China project (Z2004-1-62013), and the Young Teacher Startup Foundation project at Lanzhou University of Technology. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
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