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Microstructure and tribological properties of laser cladded self-lubricating nickel-base composite coatings containing nano-Cu and h-BN solid lubricants
Surface & Coatings Technology 359 (2019) 485–494
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Microstructure and tribological properties of laser cladded self-lubricating nickel-base composite coatings containing nano-Cu and h-BN solid lubricants Yue Zhaoa,b, Kai Fenga,b, Chengwu Yaoa,b, Pulin Niea,b, Jian Huanga,b, Zhuguo Lia,b,c,
T
⁎
a Shanghai Key Laboratory of Materials Laser Processing and Modification, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China b Collaborative Innovation Center for Advanced Ship and Deep-Sea Exploration, Shanghai 200240, China c Shanghai Innovation Institute for Materials, Shanghai 200444, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Laser cladding Self-lubricant coating Wide temperature range Nickel-base composite Steel substrate Wear
In the present work, nickel-base composite powder (Ni60), nickel-base composite powder with the addition of hBN solid lubricants (h-BN/Ni60) and nickel-base composite powder with the addition of nano-Cu encapsulated h-BN solid lubricants (nano-Cu/h-BN/Ni60) were used as raw materials to synthesize three different coatings on Q235 steels by laser cladding. Microstructures of these coatings were analyzed by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD) and Raman spectroscopy. Tribological properties of these coatings were investigated at the temperatures from 25 °C to 600 °C. High temperature microhardness measurement was performed by Vickers micro-hardness tester. The results showed that the h-BN particles survived after laser cladding and displayed a homogeneous distribution in the nickel-base composite matrix. The encapsulation of h-BN by nano-Cu resulted in an increase of h-BN content in the coating. Although the addition of nano-Cu and h-BN led to a decrease on hardness, the nano-Cu/h-BN/Ni60 coating had the lowest friction coefficient and wear rate among the three coatings in a wide range of temperature from 25 °C to 500 °C. The mechanism of wear reduction by addition of nano-Cu encapsulated h-BN solid-lubricants was also discussed in this research.
1. Introduction With the rapid development of modern manufacturing, the high temperature components such as high temperature air foil bearing (gas bearing) and gas turbine sealing is asking for the outstanding lubricating performance and wear-resistant properties in high temperature [1–6]. Actually, the high temperature components operated from room temperature to elevated temperature require stable self-lubricating performance and wear-resistant properties at a whole range of temperature to ensure the working efficiency and reliability. Solid lubricant has been used to improve the friction and wear resistant properties of materials, especially at high service temperature when liquid lubricant cannot work effectively. Hexagonal Boron Nitride (hBN) has good thermal stability, low density, well heat conductivity and electrical insulating properties and exhibits relatively low coefficient of friction at temperatures between 200 and 500 °C [7]. L.Z. Du et al. prepared a Ni3Al/h-BN composite coating by plasma spray [8] and
found that both friction coefficient and fluctuation of the friction coefficient decreased with the increasing of h-BN. H.T. Chi et al. [9] and W. Chen et al. [10] investigated the tribological behaviors of (TiB2 + hBN)/2024Al composites and Si3N4-h-BN ceramic material respectively. They both found that h-BN can reduce the friction coefficient of the composites. S.T. Zhang et al. [11] reported that the friction coefficient of Ni/h-BN coating increased with the temperature rising from 25 °C to 100 °C and then decreased gradually as the temperature rose up to 800 °C. L. Avril et al. [12] and C. Z. Shan et al. [13] investigated the tribological performances of α-Fe (Cr)-h-BN coatings and SiC/h-BN composite coating at different temperatures. They found that h-BN can only reduce the friction coefficient of the coating in the higher temperature range (200 °C–600 °C), but at lower temperature range (25 °C–200 °C) the effective of h-BN to reduce the friction coefficient was not very outstanding. Recent work concentrated on the composite with various lubricants and aimed at reducing the friction coefficient at a wide range of
⁎ Corresponding author at: Shanghai Key Laboratory of Materials Laser Processing and Modification, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China. E-mail address: [email protected] (Z. Li).
https://doi.org/10.1016/j.surfcoat.2018.12.017 Received 12 August 2018; Received in revised form 3 December 2018; Accepted 6 December 2018 Available online 06 December 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.
Surface & Coatings Technology 359 (2019) 485–494
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temperature synergistically. X.L. Shi et al. [14] investigated the tribological performance of TiAl based self-lubricating composites with addition of Ag/Ti3SiC2/BaF2/CaF2 solid lubricants by PSP. They found that because of the synergetic lubricating effect of Ag/Ti3SiC2/BaF2/ CaF2 the friction and wear properties of the composites were improved from room temperature to 600 °C. Ag formed a tribo-film to reduce the friction coefficient at low temperatures, fluorides and oxides of Ti3SiC2 reduced the friction coefficient at high temperatures. J.H. Yuan et al. [15] prepared WC-Co-Cu-BaF2/CaF2 self-lubricating coating by plasma spray. The friction and wear properties of the coatings decreased because of the synergistic self-lubricating effect of Cu and BaF2/CaF2 solid lubricants. Cu reduced the friction coefficient at low wear temperatures, while BaF2/CaF2 played the role of reducing the friction coefficient at high wear temperatures. X.L. Shi et al. fabricated Ni3Al/WS2/ Ag/h-BN intermetallic matrix composites by spark plasma sintering [16], and they found that the composites with 15 wt% WS2/Ag/h-BN lubricants possessed lower friction coefficient while the temperature rose up from 25 °C to 800 °C because of the synergetic lubricating effects of composite lubricants. S.T. Zhang et al. [17] prepared Ni3Al/hBN/Ag composite coating by reactive sintering and investigated the synergetic lubricating action of Ag and h-BN from room temperature to 800 °C. They found that Ag can reduce the friction coefficient at lower temperature range (25 °C–200 °C), h-BN played the role of self-lubricating properties at higher temperature range (200 °C–800 °C). None of the researchers above used laser cladding method, and the prepared composites had large porosity and poor bonding with the substrate. Laser cladding can make the metallurgical bond between the coating and the substrate, and the as-prepared composites have low porosity and uniform chemical composition [18]. Because of the poor wettability between h-BN and metals, h-BN is easily burned or lost in the form of splashes during the laser irradiation, resulting in a reduced deposition efficiency of the powder. It means that most of h-BN was easy to loss during the fabricated processing. H. Yan et al. used the method of high-energy ball milling to encapsulate nano-Ni onto the surface of nano-h-BN. This method improved the wettability between nano-h-BN and nickel-base matrix during YAG laser cladding process effectively [19]. But using nano-Ni to encapsulate h-BN, the solid lubricant in the as-deposited coating was still single h-BN. Without multiple solid lubricants, the friction and wear properties of the composites in a wide range of temperature cannot be improved. Soft metal copper are used as room temperature lubricants [20] usually because of their lower strength and higher plasticity extension. J.H. Shin et al. prepared the ternary Mo-Cu-N coatings on Si wafers and AISI 304 substrates by magnetron co-sputtering [21] and found that the average friction coefficient of the Mo-Cu-N coatings decreased from 0.40 to 0.21 with the increasing content of copper at room temperature. It proved that copper can reduce the friction coefficient at room temperature effectively. Besides, researchers found that copper can also reduce the friction coefficient at higher temperature. B.P. Peterson [22] prepared Cu-Mo coatings by ion-beam depositing and found that the coating was oxidized and lots of high temperature lubricants were formed such as CuO at 530 °C, which led to a decrease of friction coefficient from 0.5 to 0.2. According to these investigations, as a kind of solid lubricant, the addition of copper can reduce the friction coefficient of the composites at a wide range of temperature. At the same time, copper has good wettability in nickel based alloys. It can be completely dissolved in the nickel matrix to form a substitution solid solution [23]. In this study, therefore, h-BN was encapsulated by nano-Cu aiming at improving the wettability between h-BN and nickel-base matrix thereby increase the content of h-BN in the laser cladded composite coating and reducing the friction coefficient of the coating at a wide range of temperature.
Table 1 The chemical composition (at.%) of Ni60 alloy. Element
Ni
Cr
B
Si
C
Fe
Contain
Bal.
14.18
14.18
7.56
3.5
≤6.6
2. Experimental procedures The Ni60, h-BN/Ni60 and nano-Cu/h-BN/Ni60 coatings were prepared on low-carbon steel Q235 (normalizing) by laser cladding. Ni60 coating without solid lubricants was denoted as C1; Ni60 coating with h-BN solid lubricant (h-BN/Ni60) and Ni60 coating with nano-Cu/h-BN solid lubricants (nano-Cu/h-BN/Ni60) were represented as C2 and C3. The chemical composite of Ni60 super alloy was shown in Table 1. The morphologies of Ni60, nano-Cu and h-BN raw powder were shown in Fig. 1(a), (b) and (c). Firstly, h-BN was mixed with nano-Cu powder (4:1 in wt%) by high energy ball milling rotating at a speed of 400 rpm for 3 h to ensure the uniformity of encapsulation of nano-Cu. Fig. 1(d) and (e) showed the morphology and distribution of Cu element of the encapsulated nano-Cu/h-BN after high energy ball milling. As shown in Fig. 1(c), the raw surface of h-BN was smooth. After high energy ball milling the surface of flaky h-BN particles were covered by granular nano-Cu particles (Fig. 1(d)). The nano-Cu encapsulated h-BN particles had the same shape with the raw h-BN particles. The EDS result in Fig. 1(e) indicated that the particles on the surface of h-BN were nanoCu particles. It can be proved that the nano-Cu particles were encapsulated on the surface of h-BN particles. Secondly, the nano-Cu encapsulated h-BN powder was blended with Ni60 super alloy powder by low speed ball milling (300 rpm for 3 h). For comparison experiment, h-BN and Ni60 powder was blended through low speed ball milling. The composition and proportion of these powders was listed in Table 2. Before laser cladding, the powder was mixed with organic binder and then the mixture was pre-coated on the substrate. The pre-coated substrate was dried at 80 °C for 2 h in a vacuum oven. The laser cladding process was conducted by a continuous wave fiber laser (Laserline LDF-8000, Laserline Germany) with a maximum power of 8 kW. The spot size of the laser beam was 19 mm × 6 mm. The coatings were cladded with the power of 3 kW and the laser scanning speed is 2 mm/s. The molten pool was protected from oxidation and contamination by argon flow. Test samples were cut from the cross section of cladded layer and were mounted, polished, etched with the solution of 2 mL H2O2 and 30 mL HCl. XRD analysis was conducted by a Shimadzu XRD-6100 diffractometer with CuKα radiation operating at a step of 0.02°. The microstructure was further investigated by scanning electron microscopy (SEM, JEM7600F, JEOL Japanese) with an EDAX energy dispersive X-ray spectroscopy (EDS). The bonding situation of h-BN in nickel-base composite matrix was investigated by transmission electron microscopy (TEM, JEM-2100F, JEOL Japanese) and high resolution transmission electron microscopy-Fourier transformation (HRTEMFFT). The micro-hardness at 25 °C, 200 °C, 400 °C and 600 °C were measured by a Vickers micro-hardness tester (HTV-PHS30, ARCHIMEDES, China) with load of 1 kg and dwelling time of 15 s. The high temperature dry sliding friction tests using pin-on-disk device (UMT-2, Bruker (CETR) Germany) were conducted on each type of coatings at 25 °C, 200 °C, 300 °C, 400 °C, 500 °C and 600 °C respectively. The pins hold the Al2O3 ceramic balls with the diameter of 9.58 mm. The disks were the coating specimens which were polished, cleaned in an ultrasonic bath and dried in air. The friction tests were conducted under the load of 30 N and the rotation speed of 50 rpm. The rotation diameter and friction time of the friction test were 5 mm and 30 min respectively. The abrasion loss and friction coefficient were measured through a confocal laser scanning microscope (OLS4000, Olympus America). 486
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Fig. 1. The powder morphology of (a) Ni60 powder, (b) nano-Cu powder (c) h-BN powder (d) nano-Cu encapsulated h-BN powder and (e) the distribution of Cu element in Fig. 1(d).
structure and matrix. EDS analysis was performed on these different morphologies, and the corresponding results were reported in Table 3. In C1, according to the results of EDS and XRD tests, the bulk phases (area 1) were chromium-rich carbides/borides such as M7C3, M23C6 and CrB. Area 2 was rich in nickel and silicon, and the morphology of the phases in this area was eutectic structure. According to the XRD result, Ni2Si and Ni3Si were the main phases in this eutectic structure. Area 3 was confirmed to contain abundant of nickel and iron and can be known as γ-(Ni, M) matrix from the structure and elemental composition. In C2, the element composition of long strip phases (area 4) was similar with the bulk phases in area 1. Thus, the long strip phases were M7C3, M23C6 and CrB hard phases [27]. The morphology and the atomic percentages of phases in area 5 and 6 were similar with that in area 2 and 3. Therefore, the phases in C2 were the same with that in C1. In C3 (Fig. 3(c)), the element composition of the long strip phases in area 7 was the same with that in C2 coating. Thus, these phases were M7C3, M23C6 and CrB hard phases. The volume fraction of the hard phases was much lower than that in other coatings. The morphology of the phases in area 8 and area 9 was eutectic structure and matrix. According to the EDS result, the phase in these areas was Ni2Si, Ni3Si and γ-(Ni, M) matrix. Besides, copper was detected in these areas. Fig. 3(d) were the microstructure of a magnified area in C3. Fig. 3(e)–(h) were the EDS results of the element distribution in white rectangular area and the chemical composition in point A in Fig. 3(d). It can be observed that nitrogen and boron elements were enrichment in the black dot phases. The distribution of copper was consistent with the distribution of nickel matrix. As well know, copper and nickel are ultimate mutual solution elements [24]. Therefore, it can be proved that the nano-Cu particles originally encapsulated on the surface of h-BN were dissolved in the nickel matrix during the cladding process to form the substituted solid solution. During the process of solidification, iron and nickel elements in the molten pool formed the matrix which was consisted of γ-(Ni, M) and FeNi3 through hypoeutectic crystallization firstly [25]. In the meantime, part of the Cr and Si elements in the molten pool were dissolved in
Table 2 The component percentage (wt%) content of three different powders. Mixture no.
Ni60
h-BN
Nano-Cu
Ni60 coating (C1) h-BN/Ni60 coating (C2) Nano-Cu/h-BN/Ni60 coating (C3)
100 98.75 93.75
– 1.25 1.25
– – 5
3. Results and discussion 3.1. Microstructure The XRD patterns of C1, C2 and C3 were shown in Fig. 2. The results revealed that γ-(Ni, M), FeNi3, M23C6, M7C3, Ni2Si, Ni3Si, CrB were the main phases of the coatings. The microstructure of C1, C2 and C3 were shown in Fig. 3. It could be observed that there were three kinds of micromorphology in each coating, namely bulk phases, eutectic
Fig. 2. XRD spectra of C1, C2 and C3. 487
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Fig. 3. Microstructure of three different coatings: (a) C1, (b) C2 and (c) C3; (d) the magnified area in C3; (e–h) EDS results of the distribution of nitrogen and copper element in white rectangular area and the chemical composition of position A in (d).
the matrix to act as a solid solution strengthening. The matrix bonded the solid lubricants and the hard phases in the composite effectively [26,27]. Then chromium boron and carbon segregated to grain boundaries to form carbides like CrB, Cr7C3 and Cr23C6. As the solidification progress, partial iron atoms in the remaining liquid metal dissolved in Cr7C3 and Cr23C6 and formed stable M7C3 and M23C6 carbides.
These in-situ formed carbides and borides were hard phases in the composite, which increased the hardness of the coating effectively. Moreover, the in-situ hard phases avoided the addition of other hard phases, simplified the cladding process and improved the wear resistance of the coating [28]. Eutectic Ni2Si and Ni3Si formed between the crystals in the final stage of solidification. The eutectic structure had 488
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The interface bonding between h-BN and γ-(Ni, M) in C3 was observed by TEM and HRTEM and the results were presented in Fig. 5. As shown in Fig. 5(a), the wafer-like boron nitride could be clearly observed. Generally, there are two types of interfacial bonding in composite: mechanical bonding and chemical bonding. Chemical bonding includes diffusion transition, chemical reaction layer, and coherent and semi-coherent lattice bonding, which has strong interface bonding. The mechanical bonding interface is weak [31]. Usually, the interface between the reinforcement and the matrix in the composite is expected to form a strong chemical bonding [32]. If the bonding force between the reinforcement and the matrix is weak, the reinforcement is easily shed from the matrix during the wear process. The dropped reinforcement will form the three-body abrasive particles between the friction pair, resulting in a sharp decrease in the wear resistance of the composites [33]. Many researchers have studied the interface bonding between reinforcement and the matrix [34–36]. However, the lubricating phase has the opposite mechanism to increase the wear properties. The lubrication mechanism is expected that the detached solid lubricants form a glaze layer between the friction pair during the wear process. The glaze layer can protect the worn surface and reduce the friction coefficient of the composites [37–40]. It can be observed from Fig. 5(b) that the interface of h-BN and γ-(Ni, M) matrix was disordered. As shown in Fig. 5(c), there was no obvious orientation relationship between h-BN and γ-(Ni, M) matrix. Therefore, the bonding situation between h-BN and matrix was weak mechanical bonding. The mechanical bonding made the h-BN easily detach from the matrix during the wear process. Due to the weak van der Waals force between the h-BN sheets [41,42], the detached h-BN is easily extended to form a glaze layer under the action of the friction pair during the wear process, thereby improving the lubrication performance of the coating.
Table 3 Elemental concentrations (at.%) of the selected areas in C1, C2, C3 from EDS analysis. Spectra
Ni
Cr
Fe
C
N
B
Si
Cu
1 2 3 4 5 6 7 8 9
3.22 28.31 60.38 1.33 39.85 64.52 5.24 42.65 56.27
40.19 5.63 8.63 33.91 9.91 1.31 49.21 5.12 3.65
1.99 24.5 12.56 2.99 15.05 16.22 5.37 18.43 12.41
12.5 13.64 5.34 10.21 11.66 8.34 9.94 15.06 12.13
– – – – – – – – –
39.65 8.24 9.16 51.56 1.42 2.81 27.13 0.99 3.51
2.45 19.68 3.93 – 21.66 6.8 3.11 16.24 4.36
– – – – – – – 1.51 7.67
good toughness in the coating, which alleviated the residual stress and increased the bonding strength between the coating and the substrate effectively [29,30]. In addition, boron and silicon elements in the composites improved the self-fluxing of the nickel-based alloy and decreased its melting point, and made the excellent surface forming of the coating [28]. As shown in Fig. 3(b) and (c), black-dot particles could be obviously observed in C2 and C3. Raman spectrum test was carried out to characterize the composition of the black dots in C3. The test results were presented in Fig. 4. Raman spectra were scanned on black dot 1 and 2 in Fig. 4(a). The characteristic peak of h-BN (1365 cm−1) was found in scanning result as presented in Fig. 4(e). Therefore, the dot 1 and dot 2 were h-BN phase. The mapping results of white square area in Fig. 4(a) presented the position of h-BN phase in this area, as the bright dots shown in Fig. 4(b–d). It was obviously that the positions of bright dots in Fig. 4(b–d) were perfectly matched with the position of black dots in Fig. 4(a). So, the black dots in the coatings were h-BN phase. The volume fraction of h-BN in C3 coating was more than that in C1, as shown in Fig. 3(b) and (c). The software (image-pro) was used to statistic the specific volume fraction of h-BN phase in C2 and C3. The results revealed that the content of h-BN phase in C3 coating was 2.5 times higher than that in C2. Nano-Cu was encapsulated on the surface of h-BN uniformly, so the reflectivity of laser irradiation was increased, which protected the h-BN from burnout effectively during the laser cladding process. At the same time, it increased the wettability between h-BN and nickel-base matrix as early mentioned. Therefore, the content of h-BN in C3 coating was higher than that in C2 coating.
3.2. High temperature micro-hardness Fig. 6 showed the micro-hardness of C1, C2 and C3 at 25 °C, 200 °C, 400 °C and 600 °C. The hardness of these coatings decreased with the increase of the temperature. It implied that the coating would be soften at high temperatures. Compared with C1, the average hardness of C2 was lower. C3 had the lowest hardness from 25 °C to 600 °C. As early mentioned, the content of solid lubricants in C1, C2 and C3 increased in turn. The volume fraction of the hard phases in C3 was the lowest but its volume fraction of h-BN was highest. This led to the lowest hardness
Fig. 4. BSE microstructure and Raman result of C3; (a) BSE microstructure the selected area's Raman mapping results of (b) dot 1and (c) 2, (d) all h-BN phase in mapping area and (e) Raman shift of h-BN phase. 489
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Fig. 5. (a) TEM micrographs and (b) HRTEM images of h-BN phase in the composite coating and (c) the corresponding FFT patterns of the selected regions in (b).
temperature range, the friction coefficient of C1 is the highest. The friction coefficient of these coatings decreased in order of C1 to C2 to C3. Many references [43, 44] have been reported that h-BN has excellent lubricating property. It can effectively reduce the friction coefficient of the coating. As early mentioned, the content of h-BN in C1, C2 and C3 increased in turn. In addition, soft metal copper has excellent lubricating property in the low temperature range and it can be oxidized to CuO [45] (observed by Raman spectrum) at high temperature range (400–600 °C) which further decreased the friction coefficient. Thus, the addition of h-BN and nano-Cu solid lubricants in Ni60 coating resulted in a lowest friction coefficient of nickel-base composite coating in the temperature range from 25 °C to 600 °C. The wear rates of C1, C2 and C3 were shown in Fig. 7(b). It can be observed that the wear rates of the three coatings increased with the increasing wear temperature. According to the high temperature hardness results, the hardness of the coatings decreased with the increasing wear temperature. Many researchers have reported that the wear rate of the coating is related to its hardness [37–40]. In general, the strength and hardness of the material are proportional. Extremely high temperature during the wear tests can significantly reduce the strength of the coatings [17]. As the wear temperature increases, the strength of the coating became lower and lower. A decrease in the strength of the coatings at high temperatures caused severe adhesive wear of the coatings during the friction process. The lower the strength of the coating, the more severe adhesion wear during the wear process. Therefore, the wear rate of the coatings increased with an increase of temperature. However, besides the hardness of the coating, lubrication effect is also helpful in improving the friction and wear properties of the coating [42]. As early mentioned, the volume fraction of solid lubricants in C3 was the highest. During the wear process, the solid
Fig. 6. The micro hardness of C1, C2 and C3 at 25 °C, 200 °C, 400 °C and 600 °C.
of C3 among all tested temperatures. It can be concluded that the addition of h-BN and nano-Cu solid lubricants resulted in a decrease on micro-hardness of nickel-base composite coatings. 3.3. Tribological properties The average friction coefficient and wear rate of C1, C2 and C3 at different temperatures was presented in Fig. 7(a) and (b). From Fig. 7(a), it can be observed that as the temperature rose, the friction coefficient of these coatings decreased gradually. In the tested
Fig. 7. (a) Average friction coefficient and (b) wear rate of C1, C2 and C3 at 25 °C, 200 °C, 300 °C, 400 °C, 500 °C, 600 °C. 490
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highest at 600 °C. Fig. 9 showed the schematic diagram of the abrasive wear and adhesion wear during the wear process. Process 1, 2, 3 and 4 showed the wear mechanism of the abrasive wear. It can be seen from process 2–4 that the ceramic ball detached the hard particles from the coating under the action of friction force, leaving some pits on the worn surface. The detached hard particles formed a three-body wear in the middle of the friction pair, leaving the grooves on the worn surface of the coating. As early discussed, abrasive wear was mainly occurred at the lower wear temperature (25 °C and 200 °C). At these wear temperatures the strength of the coating was high. The wear debris was not easily detached from the coating, so the wear rate of the coating was low. Most of the detached wear debris with high hardness was difficult to extend to the glaze layer to protect the worn surface. The hard debris formed the three-body abrasive particles between the friction pair. Therefore, the abrasive wear occurred at lower wear temperatures. Process No. 1, 5 and 6 were the wear mechanism of adhesive wear. As shown in process 5, the ceramic ball detached the matrix and solid lubricants from the coating under the adhesive force to form the wear debris. Then the wear debris reattached to the worn surface to form the glaze layer. The glaze layer had the process of formation, loss, and reformation during the wear process [47–49]. So delamination can be observed on the worn surface. As early discussed, this wear mechanism occurred in the high wear temperatures (400 °C and 600 °C). At these wear temperatures the strength and hardness of the coating was low. Matrix and solid lubricants were easily detached from the coating under the action of adhesion force. The dropped wear debris was easily extended on the worn surface to form the glaze layer during the wear process. Therefore, the adhesion wear occurred at higher wear temperatures. Raman spectra of wear debris on the worn surface of C3 at 25 °C, 200 °C, 400 °C and 600 °C were conducted to further identify the effective solid lubricants in wear debris and the results were shown in Fig. 10. No oxidation reaction occurred at 25 °C, copper and h-BN were the main solid lubricants at this temperature. When the temperature increased to 200 °C, h-BN, Cr2O3 and NiO were detected in wear debris which confirmed that chromium and nickel was oxidized to Cr2O3 and NiO during the wear process. It is generally known that oxides such as Cr2O3 and NiO are excellent high temperature lubricants [50–53]. Therefore, the lubricants on the worn surface were consisted of h-BN, Cu, Cr2O3 and NiO at 200 °C. When the wear temperature was 400 °C and 600 °C, h-BN, NiO, Cr2O3 and CuO were detected in wear debris. Copper was oxidized to CuO at temperatures above 400 °C. CuO is an excellent solid lubricant and it can reduce the friction coefficient effectively [45]. Thus, h-BN, CuO, Cr2O3 and NiO were the main solid lubricants at 400 °C and 600 °C. These oxides mentioned above were produced through following reactions [45]:
lubricants can form a dense glaze layer between the friction pair [17,46]. The glaze layer can separated the real contact between the friction pair and protect the worn surface from further contact damage, thereby decrease the wear rate of the coatings [37–40,42]. Therefore, C3 had the lowest wear rate at temperature from 25 °C to 500 °C. However, the protection effect of the glaze layer is temporary. After a period of wear, the glaze layer happens to wear off, a new glaze layer will be formed on the worn surface. The glaze layer is a dynamic process of formation, loss, and reformation during the wear process [47–49]. At 600 °C, the wear rate of the three coatings increased significantly, C3 characterized with the highest wear rate. According to the high temperature hardness results, the coating severely softened at 600 °C. It can be implied that due to the sharp decrease of the strength of C3 at 600 °C, the glaze layer was loss excessively during the weather process. The protective effect of the solid lubricants was not sufficient to offset the negative impact of hardness reduction on the wear rate. Therefore, the wear rate of C3 with the lowest hardness was the highest, and the wear rate of C1 with the highest hardness was the lowest at 600 °C. The morphologies of the worn surfaces of the three coatings after the wear test at 25 °C, 200 °C, 400 °C and 600 °C were shown in Fig. 8. At 25 °C and 200 °C (Fig. 8(a)–(f)), the detachment, grooves and hard phases (carbides) can be observed clearly on the worn surface of C1, C2 and C3. The formation of grooves was due to the fact that hard phases detached from the matrix during wear process and the fallen hard phases became the three-body abrasive grains between the friction pair. It can be observed from Fig. 7(a–f) that the number of grooves on the worn surface of C1, C2 and C3 decreased progressively. Since the lubrication effect of solid lubricants in C2 and C3, the number of grooves on the worn surface of C2 and C3 were less than C1. In summary, abrasive wear was the dominating wear mechanism of C1. Microplough wear was the main wear mechanism of C2 and C3 at 25 °C and 200 °C. When the wear temperature rose up to 400 °C, the worn surfaces of C1, C2 and C3 exhibited some adhesive wear characteristics. Some discontinuity irregular wear debris were observed on the worn surface of the coatings (Fig. 8(g–i)). The hardness of C1, C2 and C3 was declined as the wear temperature rose to 400 °C. The softening of the coatings led to a decrease in the strength of the coating, so adhesive wear occurred on the worn surface. Under the effect of adhesion, the wear of the matrix and solid lubricants led to a protruding of the hard phases and the dropped debris was distributed around the steps of hard phases and matrix. The wear debris lubricated the worn surface and prevent the hard particles from detached from the matrix. Compared to 20 °C and 200 °C, the number of grooves on the worn surfaces of the coatings at 400 °C was significantly reduced. Shallow grooves were observed in C1 and C2. The worn surface of C3 was much smoother than the other coatings and grooves disappear. This indicated that the wear mechanism of C1 and C2 was adhesion wear and abrasive wear, the wear mechanism of C3 was adhesion wear at 400 °C. When the wear temperature increased to 600 °C, the hardness of the coatings was further decreased. Obvious delamination and discontinued glaze layer can be observed on the worn surfaces of C1, C2 and C3. No traces of abrasive wear such as grooves and detachment were observed on the worn surface of C1, C2 and C3. Therefore, the main wear mechanism of the coatings at 600 °C was adhesive wear. The wear debris which consisted of solid lubricants and oxide were dropped from the matrix under the effect of adhesion wear during the wear process. The dropped wear debris reattached to the worn surface under the pressure of the friction pair and formed a glaze layer. The cross section of the glaze layer was shown in Fig. 8(j–l). It can be observed that C3 had the thickest glaze layer. The glaze layer of C3 had the best protection effect of the worn surface, so the friction coefficient of C3 at 600 °C was the lowest. However, it also implied that the adhesion effect of C3 was the most serious. As the wear process progresses, the glaze layer will continuously fall and re-formation [48–50], so the wear rate of C3 was the
∆
2Ni + O2 → 2NiO ∆
4Cr3 C2 + 13O2 → 6Cr2 O3 + 8CO ∆
2Cr7 C3 + 15O2 → 7Cr2 O3 + 3CO + 3CO2 400°C
2Cu + O2 → 2CuO
(1) (2) (3) (4)
It can be concluded that solid lubricants h-BN and nano-Cu increased the wear properties of nickel-base composite coating effectively. C3 exhibited the best friction properties among tested coatings at the temperature range from 25 °C to 500 °C among which nano-Cu played the most crucial role. Firstly, nano-Cu encapsulated h-BN effectively which increased the content of h-BN in C3. Secondly, copper and h-BN decreased the friction coefficient of nickel-base composite coating in the temperature range of 25 °C to 400 °C. Thirdly, copper was oxidized to CuO. CuO together with other solid lubricants synergistically lubricated the worn surface, resulting in a further decrease of 491
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Fig. 8. SEM results of worn surfaces of C1, C2 and C3: (a, d, g, j), were the worn surface of C1 at 25 °C, 200 °C, 400 °C, 600 °C, respectively; (b, e, h, k) were the worn surface of C2 at 25 °C, 200 °C, 400 °C, 600 °C, respectively; (c, f, i, l) were the worn surface of C3 at 25 °C, 200 °C, 400 °C, 600 °C, respectively.
coatings in the temperature range from 25 °C to 500 °C. The hardness of the nickel-base composite coatings showed a downward trend with the increasing of the wear temperature. At 600 °C, the wear rate of the coating content copper and h-BN solid lubricant was the highest. The wear mechanism of the coating without solid lubricants (C1) was abrasive wear at 25 °C and 200 °C. With the addition of copper and h-BN solid lubricants, the wear mechanisms of the other two coatings were dominated by micro-ploughing wear. At 400 °C, owning to the synergetic lubrication of h-BN, Cr2O3 and NiO, the wear mechanism of the coating with lower content of solid lubricants (C1 and C2) were a mixture of adhesion wear and abrasive wear; the coating containing CuO, h-BN, Cr2O3 and NiO solid lubricants (C3) was showed adhesion wear. When the wear temperature rose up to 600 °C, serious delamination was observed on the worn surface of all coatings due to serious adhesive wear.
friction coefficient when the wear temperature above 400 °C. 4. Conclusions In this research, the addition of h-BN and nano-Cu solid lubricants into nickel-base composite coatings had been successfully carried out by laser cladding using a mixed powder. The conclusions can be drawn as follows: γ-(Ni, M), FeNi3, M23C6, M7C3, Ni2Si, Ni3Si, CrB were the main phases in these nickel-base composite coatings. h-BN particles survived in the composites after laser cladding. The encapsulation of h-BN with nano-Cu resulted in an increase of h-BN content in the coating. With the increasing content of h-BN and nano-Cu solid lubricants, the hardness of nickel-base composite coatings decreased. However, it resulted in an improvement of the wear performance of nickel-base composite 492
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Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
Microstructure and tribological properties of laser cladded self-lubricating nickel-base composite coatings containing nano-Cu and h-BN solid lubricants Yue Zhaoa,b, Kai Fenga,b, Chengwu Yaoa,b, Pulin Niea,b, Jian Huanga,b, Zhuguo Lia,b,c,
T
⁎
a Shanghai Key Laboratory of Materials Laser Processing and Modification, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China b Collaborative Innovation Center for Advanced Ship and Deep-Sea Exploration, Shanghai 200240, China c Shanghai Innovation Institute for Materials, Shanghai 200444, China
A R T I C LE I N FO
A B S T R A C T
Keywords: Laser cladding Self-lubricant coating Wide temperature range Nickel-base composite Steel substrate Wear
In the present work, nickel-base composite powder (Ni60), nickel-base composite powder with the addition of hBN solid lubricants (h-BN/Ni60) and nickel-base composite powder with the addition of nano-Cu encapsulated h-BN solid lubricants (nano-Cu/h-BN/Ni60) were used as raw materials to synthesize three different coatings on Q235 steels by laser cladding. Microstructures of these coatings were analyzed by scanning electron microscopy (SEM), transmission electron microscopy (TEM), X-ray diffraction (XRD) and Raman spectroscopy. Tribological properties of these coatings were investigated at the temperatures from 25 °C to 600 °C. High temperature microhardness measurement was performed by Vickers micro-hardness tester. The results showed that the h-BN particles survived after laser cladding and displayed a homogeneous distribution in the nickel-base composite matrix. The encapsulation of h-BN by nano-Cu resulted in an increase of h-BN content in the coating. Although the addition of nano-Cu and h-BN led to a decrease on hardness, the nano-Cu/h-BN/Ni60 coating had the lowest friction coefficient and wear rate among the three coatings in a wide range of temperature from 25 °C to 500 °C. The mechanism of wear reduction by addition of nano-Cu encapsulated h-BN solid-lubricants was also discussed in this research.
1. Introduction With the rapid development of modern manufacturing, the high temperature components such as high temperature air foil bearing (gas bearing) and gas turbine sealing is asking for the outstanding lubricating performance and wear-resistant properties in high temperature [1–6]. Actually, the high temperature components operated from room temperature to elevated temperature require stable self-lubricating performance and wear-resistant properties at a whole range of temperature to ensure the working efficiency and reliability. Solid lubricant has been used to improve the friction and wear resistant properties of materials, especially at high service temperature when liquid lubricant cannot work effectively. Hexagonal Boron Nitride (hBN) has good thermal stability, low density, well heat conductivity and electrical insulating properties and exhibits relatively low coefficient of friction at temperatures between 200 and 500 °C [7]. L.Z. Du et al. prepared a Ni3Al/h-BN composite coating by plasma spray [8] and
found that both friction coefficient and fluctuation of the friction coefficient decreased with the increasing of h-BN. H.T. Chi et al. [9] and W. Chen et al. [10] investigated the tribological behaviors of (TiB2 + hBN)/2024Al composites and Si3N4-h-BN ceramic material respectively. They both found that h-BN can reduce the friction coefficient of the composites. S.T. Zhang et al. [11] reported that the friction coefficient of Ni/h-BN coating increased with the temperature rising from 25 °C to 100 °C and then decreased gradually as the temperature rose up to 800 °C. L. Avril et al. [12] and C. Z. Shan et al. [13] investigated the tribological performances of α-Fe (Cr)-h-BN coatings and SiC/h-BN composite coating at different temperatures. They found that h-BN can only reduce the friction coefficient of the coating in the higher temperature range (200 °C–600 °C), but at lower temperature range (25 °C–200 °C) the effective of h-BN to reduce the friction coefficient was not very outstanding. Recent work concentrated on the composite with various lubricants and aimed at reducing the friction coefficient at a wide range of
⁎ Corresponding author at: Shanghai Key Laboratory of Materials Laser Processing and Modification, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China. E-mail address: [email protected] (Z. Li).
https://doi.org/10.1016/j.surfcoat.2018.12.017 Received 12 August 2018; Received in revised form 3 December 2018; Accepted 6 December 2018 Available online 06 December 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.
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temperature synergistically. X.L. Shi et al. [14] investigated the tribological performance of TiAl based self-lubricating composites with addition of Ag/Ti3SiC2/BaF2/CaF2 solid lubricants by PSP. They found that because of the synergetic lubricating effect of Ag/Ti3SiC2/BaF2/ CaF2 the friction and wear properties of the composites were improved from room temperature to 600 °C. Ag formed a tribo-film to reduce the friction coefficient at low temperatures, fluorides and oxides of Ti3SiC2 reduced the friction coefficient at high temperatures. J.H. Yuan et al. [15] prepared WC-Co-Cu-BaF2/CaF2 self-lubricating coating by plasma spray. The friction and wear properties of the coatings decreased because of the synergistic self-lubricating effect of Cu and BaF2/CaF2 solid lubricants. Cu reduced the friction coefficient at low wear temperatures, while BaF2/CaF2 played the role of reducing the friction coefficient at high wear temperatures. X.L. Shi et al. fabricated Ni3Al/WS2/ Ag/h-BN intermetallic matrix composites by spark plasma sintering [16], and they found that the composites with 15 wt% WS2/Ag/h-BN lubricants possessed lower friction coefficient while the temperature rose up from 25 °C to 800 °C because of the synergetic lubricating effects of composite lubricants. S.T. Zhang et al. [17] prepared Ni3Al/hBN/Ag composite coating by reactive sintering and investigated the synergetic lubricating action of Ag and h-BN from room temperature to 800 °C. They found that Ag can reduce the friction coefficient at lower temperature range (25 °C–200 °C), h-BN played the role of self-lubricating properties at higher temperature range (200 °C–800 °C). None of the researchers above used laser cladding method, and the prepared composites had large porosity and poor bonding with the substrate. Laser cladding can make the metallurgical bond between the coating and the substrate, and the as-prepared composites have low porosity and uniform chemical composition [18]. Because of the poor wettability between h-BN and metals, h-BN is easily burned or lost in the form of splashes during the laser irradiation, resulting in a reduced deposition efficiency of the powder. It means that most of h-BN was easy to loss during the fabricated processing. H. Yan et al. used the method of high-energy ball milling to encapsulate nano-Ni onto the surface of nano-h-BN. This method improved the wettability between nano-h-BN and nickel-base matrix during YAG laser cladding process effectively [19]. But using nano-Ni to encapsulate h-BN, the solid lubricant in the as-deposited coating was still single h-BN. Without multiple solid lubricants, the friction and wear properties of the composites in a wide range of temperature cannot be improved. Soft metal copper are used as room temperature lubricants [20] usually because of their lower strength and higher plasticity extension. J.H. Shin et al. prepared the ternary Mo-Cu-N coatings on Si wafers and AISI 304 substrates by magnetron co-sputtering [21] and found that the average friction coefficient of the Mo-Cu-N coatings decreased from 0.40 to 0.21 with the increasing content of copper at room temperature. It proved that copper can reduce the friction coefficient at room temperature effectively. Besides, researchers found that copper can also reduce the friction coefficient at higher temperature. B.P. Peterson [22] prepared Cu-Mo coatings by ion-beam depositing and found that the coating was oxidized and lots of high temperature lubricants were formed such as CuO at 530 °C, which led to a decrease of friction coefficient from 0.5 to 0.2. According to these investigations, as a kind of solid lubricant, the addition of copper can reduce the friction coefficient of the composites at a wide range of temperature. At the same time, copper has good wettability in nickel based alloys. It can be completely dissolved in the nickel matrix to form a substitution solid solution [23]. In this study, therefore, h-BN was encapsulated by nano-Cu aiming at improving the wettability between h-BN and nickel-base matrix thereby increase the content of h-BN in the laser cladded composite coating and reducing the friction coefficient of the coating at a wide range of temperature.
Table 1 The chemical composition (at.%) of Ni60 alloy. Element
Ni
Cr
B
Si
C
Fe
Contain
Bal.
14.18
14.18
7.56
3.5
≤6.6
2. Experimental procedures The Ni60, h-BN/Ni60 and nano-Cu/h-BN/Ni60 coatings were prepared on low-carbon steel Q235 (normalizing) by laser cladding. Ni60 coating without solid lubricants was denoted as C1; Ni60 coating with h-BN solid lubricant (h-BN/Ni60) and Ni60 coating with nano-Cu/h-BN solid lubricants (nano-Cu/h-BN/Ni60) were represented as C2 and C3. The chemical composite of Ni60 super alloy was shown in Table 1. The morphologies of Ni60, nano-Cu and h-BN raw powder were shown in Fig. 1(a), (b) and (c). Firstly, h-BN was mixed with nano-Cu powder (4:1 in wt%) by high energy ball milling rotating at a speed of 400 rpm for 3 h to ensure the uniformity of encapsulation of nano-Cu. Fig. 1(d) and (e) showed the morphology and distribution of Cu element of the encapsulated nano-Cu/h-BN after high energy ball milling. As shown in Fig. 1(c), the raw surface of h-BN was smooth. After high energy ball milling the surface of flaky h-BN particles were covered by granular nano-Cu particles (Fig. 1(d)). The nano-Cu encapsulated h-BN particles had the same shape with the raw h-BN particles. The EDS result in Fig. 1(e) indicated that the particles on the surface of h-BN were nanoCu particles. It can be proved that the nano-Cu particles were encapsulated on the surface of h-BN particles. Secondly, the nano-Cu encapsulated h-BN powder was blended with Ni60 super alloy powder by low speed ball milling (300 rpm for 3 h). For comparison experiment, h-BN and Ni60 powder was blended through low speed ball milling. The composition and proportion of these powders was listed in Table 2. Before laser cladding, the powder was mixed with organic binder and then the mixture was pre-coated on the substrate. The pre-coated substrate was dried at 80 °C for 2 h in a vacuum oven. The laser cladding process was conducted by a continuous wave fiber laser (Laserline LDF-8000, Laserline Germany) with a maximum power of 8 kW. The spot size of the laser beam was 19 mm × 6 mm. The coatings were cladded with the power of 3 kW and the laser scanning speed is 2 mm/s. The molten pool was protected from oxidation and contamination by argon flow. Test samples were cut from the cross section of cladded layer and were mounted, polished, etched with the solution of 2 mL H2O2 and 30 mL HCl. XRD analysis was conducted by a Shimadzu XRD-6100 diffractometer with CuKα radiation operating at a step of 0.02°. The microstructure was further investigated by scanning electron microscopy (SEM, JEM7600F, JEOL Japanese) with an EDAX energy dispersive X-ray spectroscopy (EDS). The bonding situation of h-BN in nickel-base composite matrix was investigated by transmission electron microscopy (TEM, JEM-2100F, JEOL Japanese) and high resolution transmission electron microscopy-Fourier transformation (HRTEMFFT). The micro-hardness at 25 °C, 200 °C, 400 °C and 600 °C were measured by a Vickers micro-hardness tester (HTV-PHS30, ARCHIMEDES, China) with load of 1 kg and dwelling time of 15 s. The high temperature dry sliding friction tests using pin-on-disk device (UMT-2, Bruker (CETR) Germany) were conducted on each type of coatings at 25 °C, 200 °C, 300 °C, 400 °C, 500 °C and 600 °C respectively. The pins hold the Al2O3 ceramic balls with the diameter of 9.58 mm. The disks were the coating specimens which were polished, cleaned in an ultrasonic bath and dried in air. The friction tests were conducted under the load of 30 N and the rotation speed of 50 rpm. The rotation diameter and friction time of the friction test were 5 mm and 30 min respectively. The abrasion loss and friction coefficient were measured through a confocal laser scanning microscope (OLS4000, Olympus America). 486
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Fig. 1. The powder morphology of (a) Ni60 powder, (b) nano-Cu powder (c) h-BN powder (d) nano-Cu encapsulated h-BN powder and (e) the distribution of Cu element in Fig. 1(d).
structure and matrix. EDS analysis was performed on these different morphologies, and the corresponding results were reported in Table 3. In C1, according to the results of EDS and XRD tests, the bulk phases (area 1) were chromium-rich carbides/borides such as M7C3, M23C6 and CrB. Area 2 was rich in nickel and silicon, and the morphology of the phases in this area was eutectic structure. According to the XRD result, Ni2Si and Ni3Si were the main phases in this eutectic structure. Area 3 was confirmed to contain abundant of nickel and iron and can be known as γ-(Ni, M) matrix from the structure and elemental composition. In C2, the element composition of long strip phases (area 4) was similar with the bulk phases in area 1. Thus, the long strip phases were M7C3, M23C6 and CrB hard phases [27]. The morphology and the atomic percentages of phases in area 5 and 6 were similar with that in area 2 and 3. Therefore, the phases in C2 were the same with that in C1. In C3 (Fig. 3(c)), the element composition of the long strip phases in area 7 was the same with that in C2 coating. Thus, these phases were M7C3, M23C6 and CrB hard phases. The volume fraction of the hard phases was much lower than that in other coatings. The morphology of the phases in area 8 and area 9 was eutectic structure and matrix. According to the EDS result, the phase in these areas was Ni2Si, Ni3Si and γ-(Ni, M) matrix. Besides, copper was detected in these areas. Fig. 3(d) were the microstructure of a magnified area in C3. Fig. 3(e)–(h) were the EDS results of the element distribution in white rectangular area and the chemical composition in point A in Fig. 3(d). It can be observed that nitrogen and boron elements were enrichment in the black dot phases. The distribution of copper was consistent with the distribution of nickel matrix. As well know, copper and nickel are ultimate mutual solution elements [24]. Therefore, it can be proved that the nano-Cu particles originally encapsulated on the surface of h-BN were dissolved in the nickel matrix during the cladding process to form the substituted solid solution. During the process of solidification, iron and nickel elements in the molten pool formed the matrix which was consisted of γ-(Ni, M) and FeNi3 through hypoeutectic crystallization firstly [25]. In the meantime, part of the Cr and Si elements in the molten pool were dissolved in
Table 2 The component percentage (wt%) content of three different powders. Mixture no.
Ni60
h-BN
Nano-Cu
Ni60 coating (C1) h-BN/Ni60 coating (C2) Nano-Cu/h-BN/Ni60 coating (C3)
100 98.75 93.75
– 1.25 1.25
– – 5
3. Results and discussion 3.1. Microstructure The XRD patterns of C1, C2 and C3 were shown in Fig. 2. The results revealed that γ-(Ni, M), FeNi3, M23C6, M7C3, Ni2Si, Ni3Si, CrB were the main phases of the coatings. The microstructure of C1, C2 and C3 were shown in Fig. 3. It could be observed that there were three kinds of micromorphology in each coating, namely bulk phases, eutectic
Fig. 2. XRD spectra of C1, C2 and C3. 487
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Fig. 3. Microstructure of three different coatings: (a) C1, (b) C2 and (c) C3; (d) the magnified area in C3; (e–h) EDS results of the distribution of nitrogen and copper element in white rectangular area and the chemical composition of position A in (d).
the matrix to act as a solid solution strengthening. The matrix bonded the solid lubricants and the hard phases in the composite effectively [26,27]. Then chromium boron and carbon segregated to grain boundaries to form carbides like CrB, Cr7C3 and Cr23C6. As the solidification progress, partial iron atoms in the remaining liquid metal dissolved in Cr7C3 and Cr23C6 and formed stable M7C3 and M23C6 carbides.
These in-situ formed carbides and borides were hard phases in the composite, which increased the hardness of the coating effectively. Moreover, the in-situ hard phases avoided the addition of other hard phases, simplified the cladding process and improved the wear resistance of the coating [28]. Eutectic Ni2Si and Ni3Si formed between the crystals in the final stage of solidification. The eutectic structure had 488
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The interface bonding between h-BN and γ-(Ni, M) in C3 was observed by TEM and HRTEM and the results were presented in Fig. 5. As shown in Fig. 5(a), the wafer-like boron nitride could be clearly observed. Generally, there are two types of interfacial bonding in composite: mechanical bonding and chemical bonding. Chemical bonding includes diffusion transition, chemical reaction layer, and coherent and semi-coherent lattice bonding, which has strong interface bonding. The mechanical bonding interface is weak [31]. Usually, the interface between the reinforcement and the matrix in the composite is expected to form a strong chemical bonding [32]. If the bonding force between the reinforcement and the matrix is weak, the reinforcement is easily shed from the matrix during the wear process. The dropped reinforcement will form the three-body abrasive particles between the friction pair, resulting in a sharp decrease in the wear resistance of the composites [33]. Many researchers have studied the interface bonding between reinforcement and the matrix [34–36]. However, the lubricating phase has the opposite mechanism to increase the wear properties. The lubrication mechanism is expected that the detached solid lubricants form a glaze layer between the friction pair during the wear process. The glaze layer can protect the worn surface and reduce the friction coefficient of the composites [37–40]. It can be observed from Fig. 5(b) that the interface of h-BN and γ-(Ni, M) matrix was disordered. As shown in Fig. 5(c), there was no obvious orientation relationship between h-BN and γ-(Ni, M) matrix. Therefore, the bonding situation between h-BN and matrix was weak mechanical bonding. The mechanical bonding made the h-BN easily detach from the matrix during the wear process. Due to the weak van der Waals force between the h-BN sheets [41,42], the detached h-BN is easily extended to form a glaze layer under the action of the friction pair during the wear process, thereby improving the lubrication performance of the coating.
Table 3 Elemental concentrations (at.%) of the selected areas in C1, C2, C3 from EDS analysis. Spectra
Ni
Cr
Fe
C
N
B
Si
Cu
1 2 3 4 5 6 7 8 9
3.22 28.31 60.38 1.33 39.85 64.52 5.24 42.65 56.27
40.19 5.63 8.63 33.91 9.91 1.31 49.21 5.12 3.65
1.99 24.5 12.56 2.99 15.05 16.22 5.37 18.43 12.41
12.5 13.64 5.34 10.21 11.66 8.34 9.94 15.06 12.13
– – – – – – – – –
39.65 8.24 9.16 51.56 1.42 2.81 27.13 0.99 3.51
2.45 19.68 3.93 – 21.66 6.8 3.11 16.24 4.36
– – – – – – – 1.51 7.67
good toughness in the coating, which alleviated the residual stress and increased the bonding strength between the coating and the substrate effectively [29,30]. In addition, boron and silicon elements in the composites improved the self-fluxing of the nickel-based alloy and decreased its melting point, and made the excellent surface forming of the coating [28]. As shown in Fig. 3(b) and (c), black-dot particles could be obviously observed in C2 and C3. Raman spectrum test was carried out to characterize the composition of the black dots in C3. The test results were presented in Fig. 4. Raman spectra were scanned on black dot 1 and 2 in Fig. 4(a). The characteristic peak of h-BN (1365 cm−1) was found in scanning result as presented in Fig. 4(e). Therefore, the dot 1 and dot 2 were h-BN phase. The mapping results of white square area in Fig. 4(a) presented the position of h-BN phase in this area, as the bright dots shown in Fig. 4(b–d). It was obviously that the positions of bright dots in Fig. 4(b–d) were perfectly matched with the position of black dots in Fig. 4(a). So, the black dots in the coatings were h-BN phase. The volume fraction of h-BN in C3 coating was more than that in C1, as shown in Fig. 3(b) and (c). The software (image-pro) was used to statistic the specific volume fraction of h-BN phase in C2 and C3. The results revealed that the content of h-BN phase in C3 coating was 2.5 times higher than that in C2. Nano-Cu was encapsulated on the surface of h-BN uniformly, so the reflectivity of laser irradiation was increased, which protected the h-BN from burnout effectively during the laser cladding process. At the same time, it increased the wettability between h-BN and nickel-base matrix as early mentioned. Therefore, the content of h-BN in C3 coating was higher than that in C2 coating.
3.2. High temperature micro-hardness Fig. 6 showed the micro-hardness of C1, C2 and C3 at 25 °C, 200 °C, 400 °C and 600 °C. The hardness of these coatings decreased with the increase of the temperature. It implied that the coating would be soften at high temperatures. Compared with C1, the average hardness of C2 was lower. C3 had the lowest hardness from 25 °C to 600 °C. As early mentioned, the content of solid lubricants in C1, C2 and C3 increased in turn. The volume fraction of the hard phases in C3 was the lowest but its volume fraction of h-BN was highest. This led to the lowest hardness
Fig. 4. BSE microstructure and Raman result of C3; (a) BSE microstructure the selected area's Raman mapping results of (b) dot 1and (c) 2, (d) all h-BN phase in mapping area and (e) Raman shift of h-BN phase. 489
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Fig. 5. (a) TEM micrographs and (b) HRTEM images of h-BN phase in the composite coating and (c) the corresponding FFT patterns of the selected regions in (b).
temperature range, the friction coefficient of C1 is the highest. The friction coefficient of these coatings decreased in order of C1 to C2 to C3. Many references [43, 44] have been reported that h-BN has excellent lubricating property. It can effectively reduce the friction coefficient of the coating. As early mentioned, the content of h-BN in C1, C2 and C3 increased in turn. In addition, soft metal copper has excellent lubricating property in the low temperature range and it can be oxidized to CuO [45] (observed by Raman spectrum) at high temperature range (400–600 °C) which further decreased the friction coefficient. Thus, the addition of h-BN and nano-Cu solid lubricants in Ni60 coating resulted in a lowest friction coefficient of nickel-base composite coating in the temperature range from 25 °C to 600 °C. The wear rates of C1, C2 and C3 were shown in Fig. 7(b). It can be observed that the wear rates of the three coatings increased with the increasing wear temperature. According to the high temperature hardness results, the hardness of the coatings decreased with the increasing wear temperature. Many researchers have reported that the wear rate of the coating is related to its hardness [37–40]. In general, the strength and hardness of the material are proportional. Extremely high temperature during the wear tests can significantly reduce the strength of the coatings [17]. As the wear temperature increases, the strength of the coating became lower and lower. A decrease in the strength of the coatings at high temperatures caused severe adhesive wear of the coatings during the friction process. The lower the strength of the coating, the more severe adhesion wear during the wear process. Therefore, the wear rate of the coatings increased with an increase of temperature. However, besides the hardness of the coating, lubrication effect is also helpful in improving the friction and wear properties of the coating [42]. As early mentioned, the volume fraction of solid lubricants in C3 was the highest. During the wear process, the solid
Fig. 6. The micro hardness of C1, C2 and C3 at 25 °C, 200 °C, 400 °C and 600 °C.
of C3 among all tested temperatures. It can be concluded that the addition of h-BN and nano-Cu solid lubricants resulted in a decrease on micro-hardness of nickel-base composite coatings. 3.3. Tribological properties The average friction coefficient and wear rate of C1, C2 and C3 at different temperatures was presented in Fig. 7(a) and (b). From Fig. 7(a), it can be observed that as the temperature rose, the friction coefficient of these coatings decreased gradually. In the tested
Fig. 7. (a) Average friction coefficient and (b) wear rate of C1, C2 and C3 at 25 °C, 200 °C, 300 °C, 400 °C, 500 °C, 600 °C. 490
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highest at 600 °C. Fig. 9 showed the schematic diagram of the abrasive wear and adhesion wear during the wear process. Process 1, 2, 3 and 4 showed the wear mechanism of the abrasive wear. It can be seen from process 2–4 that the ceramic ball detached the hard particles from the coating under the action of friction force, leaving some pits on the worn surface. The detached hard particles formed a three-body wear in the middle of the friction pair, leaving the grooves on the worn surface of the coating. As early discussed, abrasive wear was mainly occurred at the lower wear temperature (25 °C and 200 °C). At these wear temperatures the strength of the coating was high. The wear debris was not easily detached from the coating, so the wear rate of the coating was low. Most of the detached wear debris with high hardness was difficult to extend to the glaze layer to protect the worn surface. The hard debris formed the three-body abrasive particles between the friction pair. Therefore, the abrasive wear occurred at lower wear temperatures. Process No. 1, 5 and 6 were the wear mechanism of adhesive wear. As shown in process 5, the ceramic ball detached the matrix and solid lubricants from the coating under the adhesive force to form the wear debris. Then the wear debris reattached to the worn surface to form the glaze layer. The glaze layer had the process of formation, loss, and reformation during the wear process [47–49]. So delamination can be observed on the worn surface. As early discussed, this wear mechanism occurred in the high wear temperatures (400 °C and 600 °C). At these wear temperatures the strength and hardness of the coating was low. Matrix and solid lubricants were easily detached from the coating under the action of adhesion force. The dropped wear debris was easily extended on the worn surface to form the glaze layer during the wear process. Therefore, the adhesion wear occurred at higher wear temperatures. Raman spectra of wear debris on the worn surface of C3 at 25 °C, 200 °C, 400 °C and 600 °C were conducted to further identify the effective solid lubricants in wear debris and the results were shown in Fig. 10. No oxidation reaction occurred at 25 °C, copper and h-BN were the main solid lubricants at this temperature. When the temperature increased to 200 °C, h-BN, Cr2O3 and NiO were detected in wear debris which confirmed that chromium and nickel was oxidized to Cr2O3 and NiO during the wear process. It is generally known that oxides such as Cr2O3 and NiO are excellent high temperature lubricants [50–53]. Therefore, the lubricants on the worn surface were consisted of h-BN, Cu, Cr2O3 and NiO at 200 °C. When the wear temperature was 400 °C and 600 °C, h-BN, NiO, Cr2O3 and CuO were detected in wear debris. Copper was oxidized to CuO at temperatures above 400 °C. CuO is an excellent solid lubricant and it can reduce the friction coefficient effectively [45]. Thus, h-BN, CuO, Cr2O3 and NiO were the main solid lubricants at 400 °C and 600 °C. These oxides mentioned above were produced through following reactions [45]:
lubricants can form a dense glaze layer between the friction pair [17,46]. The glaze layer can separated the real contact between the friction pair and protect the worn surface from further contact damage, thereby decrease the wear rate of the coatings [37–40,42]. Therefore, C3 had the lowest wear rate at temperature from 25 °C to 500 °C. However, the protection effect of the glaze layer is temporary. After a period of wear, the glaze layer happens to wear off, a new glaze layer will be formed on the worn surface. The glaze layer is a dynamic process of formation, loss, and reformation during the wear process [47–49]. At 600 °C, the wear rate of the three coatings increased significantly, C3 characterized with the highest wear rate. According to the high temperature hardness results, the coating severely softened at 600 °C. It can be implied that due to the sharp decrease of the strength of C3 at 600 °C, the glaze layer was loss excessively during the weather process. The protective effect of the solid lubricants was not sufficient to offset the negative impact of hardness reduction on the wear rate. Therefore, the wear rate of C3 with the lowest hardness was the highest, and the wear rate of C1 with the highest hardness was the lowest at 600 °C. The morphologies of the worn surfaces of the three coatings after the wear test at 25 °C, 200 °C, 400 °C and 600 °C were shown in Fig. 8. At 25 °C and 200 °C (Fig. 8(a)–(f)), the detachment, grooves and hard phases (carbides) can be observed clearly on the worn surface of C1, C2 and C3. The formation of grooves was due to the fact that hard phases detached from the matrix during wear process and the fallen hard phases became the three-body abrasive grains between the friction pair. It can be observed from Fig. 7(a–f) that the number of grooves on the worn surface of C1, C2 and C3 decreased progressively. Since the lubrication effect of solid lubricants in C2 and C3, the number of grooves on the worn surface of C2 and C3 were less than C1. In summary, abrasive wear was the dominating wear mechanism of C1. Microplough wear was the main wear mechanism of C2 and C3 at 25 °C and 200 °C. When the wear temperature rose up to 400 °C, the worn surfaces of C1, C2 and C3 exhibited some adhesive wear characteristics. Some discontinuity irregular wear debris were observed on the worn surface of the coatings (Fig. 8(g–i)). The hardness of C1, C2 and C3 was declined as the wear temperature rose to 400 °C. The softening of the coatings led to a decrease in the strength of the coating, so adhesive wear occurred on the worn surface. Under the effect of adhesion, the wear of the matrix and solid lubricants led to a protruding of the hard phases and the dropped debris was distributed around the steps of hard phases and matrix. The wear debris lubricated the worn surface and prevent the hard particles from detached from the matrix. Compared to 20 °C and 200 °C, the number of grooves on the worn surfaces of the coatings at 400 °C was significantly reduced. Shallow grooves were observed in C1 and C2. The worn surface of C3 was much smoother than the other coatings and grooves disappear. This indicated that the wear mechanism of C1 and C2 was adhesion wear and abrasive wear, the wear mechanism of C3 was adhesion wear at 400 °C. When the wear temperature increased to 600 °C, the hardness of the coatings was further decreased. Obvious delamination and discontinued glaze layer can be observed on the worn surfaces of C1, C2 and C3. No traces of abrasive wear such as grooves and detachment were observed on the worn surface of C1, C2 and C3. Therefore, the main wear mechanism of the coatings at 600 °C was adhesive wear. The wear debris which consisted of solid lubricants and oxide were dropped from the matrix under the effect of adhesion wear during the wear process. The dropped wear debris reattached to the worn surface under the pressure of the friction pair and formed a glaze layer. The cross section of the glaze layer was shown in Fig. 8(j–l). It can be observed that C3 had the thickest glaze layer. The glaze layer of C3 had the best protection effect of the worn surface, so the friction coefficient of C3 at 600 °C was the lowest. However, it also implied that the adhesion effect of C3 was the most serious. As the wear process progresses, the glaze layer will continuously fall and re-formation [48–50], so the wear rate of C3 was the
∆
2Ni + O2 → 2NiO ∆
4Cr3 C2 + 13O2 → 6Cr2 O3 + 8CO ∆
2Cr7 C3 + 15O2 → 7Cr2 O3 + 3CO + 3CO2 400°C
2Cu + O2 → 2CuO
(1) (2) (3) (4)
It can be concluded that solid lubricants h-BN and nano-Cu increased the wear properties of nickel-base composite coating effectively. C3 exhibited the best friction properties among tested coatings at the temperature range from 25 °C to 500 °C among which nano-Cu played the most crucial role. Firstly, nano-Cu encapsulated h-BN effectively which increased the content of h-BN in C3. Secondly, copper and h-BN decreased the friction coefficient of nickel-base composite coating in the temperature range of 25 °C to 400 °C. Thirdly, copper was oxidized to CuO. CuO together with other solid lubricants synergistically lubricated the worn surface, resulting in a further decrease of 491
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Fig. 8. SEM results of worn surfaces of C1, C2 and C3: (a, d, g, j), were the worn surface of C1 at 25 °C, 200 °C, 400 °C, 600 °C, respectively; (b, e, h, k) were the worn surface of C2 at 25 °C, 200 °C, 400 °C, 600 °C, respectively; (c, f, i, l) were the worn surface of C3 at 25 °C, 200 °C, 400 °C, 600 °C, respectively.
coatings in the temperature range from 25 °C to 500 °C. The hardness of the nickel-base composite coatings showed a downward trend with the increasing of the wear temperature. At 600 °C, the wear rate of the coating content copper and h-BN solid lubricant was the highest. The wear mechanism of the coating without solid lubricants (C1) was abrasive wear at 25 °C and 200 °C. With the addition of copper and h-BN solid lubricants, the wear mechanisms of the other two coatings were dominated by micro-ploughing wear. At 400 °C, owning to the synergetic lubrication of h-BN, Cr2O3 and NiO, the wear mechanism of the coating with lower content of solid lubricants (C1 and C2) were a mixture of adhesion wear and abrasive wear; the coating containing CuO, h-BN, Cr2O3 and NiO solid lubricants (C3) was showed adhesion wear. When the wear temperature rose up to 600 °C, serious delamination was observed on the worn surface of all coatings due to serious adhesive wear.
friction coefficient when the wear temperature above 400 °C. 4. Conclusions In this research, the addition of h-BN and nano-Cu solid lubricants into nickel-base composite coatings had been successfully carried out by laser cladding using a mixed powder. The conclusions can be drawn as follows: γ-(Ni, M), FeNi3, M23C6, M7C3, Ni2Si, Ni3Si, CrB were the main phases in these nickel-base composite coatings. h-BN particles survived in the composites after laser cladding. The encapsulation of h-BN with nano-Cu resulted in an increase of h-BN content in the coating. With the increasing content of h-BN and nano-Cu solid lubricants, the hardness of nickel-base composite coatings decreased. However, it resulted in an improvement of the wear performance of nickel-base composite 492
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Fig. 10. Raman shift of wear debris on C3.
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