- Email: [email protected]
Abrasion resistance of Ni-Cr-B-Si coating deposited by laser cladding process
Tribology International 143 (2020) 106002
Contents lists available at ScienceDirect
Tribology International journal homepage: http://www.elsevier.com/locate/triboint
Abrasion resistance of Ni-Cr-B-Si coating deposited by laser cladding process☆ � de Sousa a, *, Francisco Ratusznei a, Milton Pereira a, Jurandir Marcos Sa Richard de Medeiros Castro b, Elvys Isaías Mercado Curi b a b
Laborat� orio de Mec^ anica de Precis~ ao, Universidade Federal de Santa Catarina, Florian� opolis, Brazil Laborat� orio de Vibraç~ oes e Tribologia, Faculdade SATC, Criciúma, Brazil
A R T I C L E I N F O
A B S T R A C T
Keywords: Abrasion resistance Chromium carbides Thermal gradient Cooling cracks
Ni-Cr-B-Si alloy deposited by laser cladding has its tribological performance evaluated. Deposition parameters selection is a difficult task. In this work, Ni-Cr-B-Si coatings were deposited by a fiber laser source with a coaxial powder nozzle on low carbon steel substrate. Microstructure, microhardness and abrasive wear tests were per formed. Results showed coatings with good surface adhesion. Microhardness was 10% higher in coatings with higher concentration and phase size of chromium carbides. Volumetric loss and wear coefficient showed 10% variation between the coatings. Coatings with different thermal gradients showed different dilution levels, affecting their abrasion resistance. The better tribological performance was obtained for coatings with lower cooling cracks density and higher volumetric fraction of carbides, which both mitigated wear micromechanism action.
1. Introduction Currently, surface engineering has proposed new technologies and alloys to enhance the performance of mechanical components. The materials subjected to severe wear require coatings that increase the lifetime of machine elements, and consequently, minimize financial losses due to shutdowns for the maintenance and replacement of worn components. Among the various types, abrasive wear mechanisms are the most frequently encountered, resulting from the friction and sliding between the hard particles and the surface of solid components. Industry accounts for approximately 30% of total world energy consumption [1]. About 23% of these 30% come from tribological contacts [2]. Sectors of great economic importance, such as the agri cultural, oil and gas and mining industry, requires and are interested in increasing their components lifetime. In mining, for example, 40% of the expenses related to friction losses, 27% of the deteriorated parts replacement, 26% of the workforce and 7% of production losses can be avoided using efficient tribological solutions [3]. The environment can also be greatly benefited by advanced and efficient tribological solutions implementation. An extensive study has shown that in the short term (approximately 8 years), on a global scale,
due to the implementation of advanced tribological technologies, CO2 emission reduction can be as much as 1.460 Mt, representing a saving of 455.000 € (Euros). In the long term (around 15 years), these values rise to 3.140 Mt and 970.000 €, respectively [2]. According to this context, it is important to improve the tribological properties, especially abrasive wear resistance, which causes volume loss, dimensional changes and failures in mechanical and related com ponents. Thus, the search for alternative manufacturing processes that provide cost-effective features has become a constant aim for research centers around the world. One adopted solution is the application of hard coatings by deposi tion processes, such as thermal spraying and arc welding processes. In these techniques, layers with special properties are deposited over a lower mechanical strength substrate. Instead of manufacturing parts composed entirely by special materials, only the surface area of interest is coated, and this approach can drastically reduce the components final cost. Despite the cited processes can generate acceptable results, there are still challenges to be overcome to allow high quality coatings production without problems related to distortions appearance, lack of metallur gical adhesion between substrate and addition material, extensive heat
This Work was presented at TRIBOBR2018. * Corresponding author. E-mail address: [email protected] (J.M.S. Sousa).
☆
https://doi.org/10.1016/j.triboint.2019.106002 Received 4 August 2019; Received in revised form 1 October 2019; Accepted 3 October 2019 Available online 16 October 2019 0301-679X/© 2019 Elsevier Ltd. All rights reserved.
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
affected zone (HAZ) and high dilution level [4–7]. In order to solve these problems commonly related to conventional arc and spraying processes, several studies have been carried out involving other solutions. Laser cladding technique, for instance, has the ability to overcome these challenges, generating better coatings when compared to other deposition methods. A number of these features is provided by the unique properties from laser radiation, which are: high directionality beam, high power available, high versatility and achiev able energy density, which allow the processing of a wide variety of materials and alloys, including the processing of complex geometry components [4,6]. Nickel (Ni) based alloys are characterized by having good abrasion and corrosion resistance properties, even in high-temperature environ ments. Among these, Ni-Cr-B-Si class alloys, which have good resistance to the mentioned wear mechanisms, also presents high hardness values [8–10]. These alloys are classified as self-fluxing, a term referring to alloys capable of reducing their liquidus temperature, allowing the hard phases precipitation. Each constituent element of these alloys has a specific purpose: Cr (chromium) protects against corrosion at high temperatures and inserts hard phases; B (boron) reduces the melting point and also contributes to the hard phases increase; Si (silicon) is responsible for self-fluxing properties and C (carbon) inserts wear-resistant carbides [11–13]. Ni-based alloys are widely applied in components from various in dustries, such as the oil and gas industry, glass molds, cold metalwork components and slurry purge elements. Besides to the cited examples, Ni-Cr-B-Si class alloys are highly suitable for maintenance and repair in industrial environments, such as extrusion screws for the plastic in dustry [13,14]. Such alloys are commonly deposited by thermal spraying processes. However, besides to the previously mentioned problems for these techniques, coating layers with high porosity concentration are also generated [11]. To fill this gap, laser cladding process is being applied, which provides the desired characteristic through a highly concentrated beam, mitigating the porosity while inserts less heat on the substrate. Laser cladding characteristics, combined with the choice of materials with good mechanical properties and suitable processing parameters, allow the high-quality coatings production. However, parametrizing the process adequately is a complex task, since there are an extensive number of variables involved, making the process multidisciplinary, requiring knowledge in the physics and metallurgy of the deposition process and also materials science. This paper aims to contribute to this discussion by evaluating the abrasion resistance of Ni-Cr-B-Si alloy, deposited by laser cladding with different processing parameters levels on AISI 1020 carbon steel sub strate, applying the wear abrasion test by using the standard ASTM G65. The processed coatings will be subjected to dilution, microhardness, microstructure, volumetric loss, Archard wear coefficient and worn surfaces aspects analysis. The main objectives are to evaluate the deposition variables effects, the characteristics and properties of the alloy on the coatings obtained quality, besides establishing a processing methodology for the manufacture of coatings with this deposition technique and alloy.
(76.0 � 25.4 � 10.0 mm) were prepared, with average roughness Ra ¼ 0.7 μm (test surface). Ni-Cr-B-Si alloy 1545-00 (powder) from €gana €s S.A manufacturer was used as the coating material. The alloy is Ho suitable for coatings, has a hardness value of 490 HV30 and a density of 7.1 g/cm3. The particle powder size of 150.0 þ 53.0 μm, with rounded morphology, was observed by scanning electron microscopy (SEM). Due to the laser cladding nozzle and powder feeder requirements, alloy was sifted to the granulometric range of 106 þ 53.0 μm. Table 1 shows its chemical composition. 2.2. Coatings deposition process Prior to the deposition process, substrate samples were subjected to an abrasive blasting process (CMV GS-9075X equipment with SINTER BALL NORBLAST sand) in order to reach an opaque surface, avoiding reflectivity, to reach a uniform material absorptivity during processing. Thereafter, they were cleaned in ethyl alcohol to ensure impurities removal. Samples were coated using an Yb-Ytterbium fiber laser source from the IPG PHOTONICS® manufacturer, with maximum power of 10-kW, wavelength (λ) ¼ between 1070 nm and 1080 nm, laser beam diameter at beam waist ¼ 0.85 mm, M2 ¼ 1.1. Specimens are moved using a numerically controlled XYZ coordinate table. Coating material is added by a disk type powder feeding system (GTV) with a1powder feed rate _ ¼ between 0.5 and 30.0 g/min. Powder feeder nozzle is the coaxial (m) type. Argon (Ar) was used as protection and carrier gas. Deposition parameters were determined based on 24 single beads preliminary tests, where2laser beam power (P) and3scanning speed (V) were varied, keeping the other parameters constant. Fig. 1 flowchart simplifies the steps of this approach. Coated specimens resulting from this step were subjected to a standard metallography procedure: sanding (80, 220, 400, 600 and 1200 Mesh), polishing with 0.3 μm alumina and chemical etching in 2.0% nital reagent for 3.0 s. These specimens were subjected to optical microscopy. Using the Image J software, single beads morphology was evaluated and the ones with lowest dilution values, without defects and/or discontinuities such cracks, pores, de tachments at the interface substrate-coating, and with better geometric aspects, were selected. The methodology considered: height (h � 0.3 mm), width (w � 1.0 mm), area, substrate penetration (p lowest possible) and wettability angle (β < less than 90� to avoid pore insertion in the overlap step). The final deposition parameter was the overlapping condition. The overlap rate was estimated from predictions provided by a code pro grammed into MATLAB software, which was developed based on the work presented by Ocelík et al. [15]. After inserting the input data (modeling parameters: clad width - mm, clad height - mm, clad overlap %, number of clad tracks and layers) into the algorithm and performing the interactions, cross sections with the best geometric surface and resulted coating height were obtained. In this way, the final deposition parameters selection was based on the relation between the following aspects: 1 - lower dilution values; 2 no cracks; 3 - absence or fewer pores presence; 4 - geometric Table 1 Ni-Cr-B-Si 1545-00 alloy chemical composition.
2. Methodology
Elements (wt.%)
2.1. Substrate and coating material AISI 1020 laminated steel with chemical composition wt.%: 0.18–0.23 (C); 0.3–0.6 (Mn); 0.04 (P); 0.05 (S) and (Fe) in balance was used as the substrate material. The choice of this material is due to its wide industrial applicability. In this way, its choice adequately simulates a real application condition. The substrate was subjected to machining process (cutting, milling and grinding) in order to satisfy the dimensions and surface quality required by ASTM G65 standard. Rectangular specimen
C
Fe
Cr
B
Si
Ni
0.3–0.4
2.0–3.4
7.5–10.0
1.7–2.0
3.3–3.9
Bal.
1 2 3
2
_ powder feed rate. m: P: laser beam power. V: scanning speed.
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 1. Deposition parameters determination.
characteristics and 5 - overlap rate prediction. 3 of the 24 parameter combinations were chosen to deposit the coatings. These are presented in Table 2, being4coatings 1 (C1), 2 (C2) and 3 (C3), respectively. These parameters shown in Table 2 were combined with fixed pa _ ¼ 10.4 g/min); laser beam diameter rameters, such: powder feed rate (m at beam waist (0.85 mm), distance from substrate surface to beam waist (35.0 mm), shielding gas flow (Ar ¼ 15.0 l/min) and carrier gas flow (Ar ¼ 5.0 l/min), to deposit the coatings. For each test group, two sam ples with rectangular area (20 � 25 mm) were deposited to be tested and analyzed (dilution, microhardness and microstructure).
coating), cross-sectional samples were evaluated by optical microscopy (OM) and scanning electron microscopy (SEM) using Backscattering Electron Detector (BSE) a HITACHI device, model TM-3030. Specific regions (presented particular or abnormal aspects) had their chemical composition evaluated by X-ray dispersive energy spectroscopy (EDX) in the SWIFTED 3000 apparatus, coupled to the SEM. 2.4. Tribological test For the tribological tests, two samples of each coating condition were prepared. In order to qualify the Ni-Cr-B-Si 1545-00 alloy abrasion resistance deposited by laser cladding, the dry sand/rubber wheel abrasion test was carried out, as stated by ASTM G65 (2016) [17]. The specimens were cleaned using ethyl alcohol and weighted in a precision scale (MARTE AD500 of resolution 0.001 g) before the start and after each test section in order to measure the volumetric loss due to tribological tests. The tribometer was previously calibrated according to the ASTM G65 requirements for sliding speed, operating stability and abrasive mass flow. The sand used as abrasive underwent a sieving process to reach a 30 Mesh grain size, followed by heating at 100 � C for 1 h in order to remove the moisture. Using ASTM G65 procedure A, the following parameters were selected: 130 N load, 200 rpm speed, 30 min test time, totaling a sliding distance of 4309 m, under a 390 g/min continuous sand flow. In order to perform an analysis of wear tendency with respect to time, the following procedure was performed: each sample was continuously tested for 5 min; after every 5 min of testing, the sample was cleaned again and weighed; after this step, the sample was reposi tioned in the tribometer. This process was performed six times for each specimen. Fig. 2 summarizes the main procedures carried out in the samples: (a) coating as deposited condition, (b) as grinded condition and (c) worn surface after tribological test.
2.3. Mechanical and microstructural characterization Mechanical characterization was carried out with the selection of one sample from each coating condition, randomly chosen. Dilution was measured in the final coatings cross-sections in a procedure analogous to performed on the single beads’ specimens. Coatings hardness was evaluated through Vickers microhardness test, following the specifications indicated by ASTM E92 (2017) [16]: parameters regarding strength, test time and a minimum distance be tween indentations, using the SHIMADZU HMV-2TADW equipment. The HV1 scale with a 10.0 N load for 10.0 s, was used. For the hardness evaluation of the coatings surface, an indentation pattern was elabo rated (in order to obtaining a reliable and low standard deviation result): 20 uniformly distributed indentations (5 mm spacing between each indentation) were made in a rectangular pattern (5 � 4) along the sample length, in order to sweep the region subsequently submitted to the tribological test. To evaluate the hardness profile, following pattern was adopted: 3 indentation lines spaced in 6 mm each other were drawn in the sample width direction (total sample width: 25 mm). The average coatings thickness was 1.5 mm, so 5 indentations were spaced 0.4 mm across the sample height, so that it was possible to reach the different areas (coating, diluted region, HAZ and substrate). Preparation process for metallographic analysis was similar to that used in the single beads preliminary tests. The only difference was the etching, which was performed by using a reagent composed of 10 ml FeCl3, 20 ml HCl, 50 ml H2O and 1 ml HNO3 for 90 s. In order to evaluate the different microstructure areas (substrate, diluted region, HAZ and
2.5. Wear analysis Results from the ASTM G65 tribological tests were evaluated based
Table 2 Coatings deposition parameters. Parameters
C1
C2
C3
Laser beam power (kW) Scanning speed (mm/s) Overlap (%) Energy imput (J/mm) Dilution (%)
1.05 5.0 30 210.0 4.1
1.40 21.7 30 64.6 12.6
1.75 30.0 30 58.3 27.4
4
Fig. 2. Main steps applied to samples: (a) as coated, (b) as grinded and (c) after tribological test.
Coating 1 (C1), coating 2 (C2) and coating 3 (C3). 3
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
on the following perspectives:
phase and also identified cooling cracks, which ceased after preheating the substrate at 400 � C and 350 � C, respectively [21,22]. Besides to the cited researches, there are other authors in the literature that highlight similar results. It is important to note that cracks can also be nucleated from defects in microstructure or inclusions in the material. In the present research, as no impurities of this nature were detected, observed cracks were conferred to the cooling phenomena. Coatings dilution was also consistent with deposition parameters employed and the single beads preliminary tests: C1 ¼ 2.6%, C2 ¼ 4.8% and C3 ¼ 22.6%. It is complex to infer about the deposition/dilution _ parameters relationship. It is likely that the result is connected to m used. As it was maintained constant, it is believed that the higher V used in C2 and C3 provided lower powder insertion per time unit and, even increasing P, an optimum condition was not reached, where the action of both three parameters generated coatings with similar geometric and metallurgical aspects. The dilution impacted the chemical composition, a factor evidenced by EDX analysis, which revealed an increase in Fe content and Ni reduction in the higher dilution coatings. In relation to Cr, B and Si, differences were not significant. In the literature, several papers describe hardness and wear resistance reduction proportional to the coating’s dilution increase, indicating the substrate chemical composition and deposition technique as key factors. Several authors emphasize that this behavior is the major responsible for hardness decreasing, a factor consistent with this research results [8,10,23]. Besides to the proper metallurgical adhesion, three microstructure characteristic regions can be visualized: a precipitate free band (planar), located near the fusion line, a hypoeutectic and a eutectic layer. These changes resulted from chemical composition variation, especially Fe content, where, in the planar and hypoeutectic (high concentration) and eutectic regions (lower content towards the coating top). Fig. 5 (ob tained via OM) indicates these different characteristic regions. In Fig. 5, dilution effect was evident through the precipitate free band establishment, which is larger (has a greater thickness) in C3. This characteristic is possible connected to variation in chemical composi tion, more specifically, the Fe content (the main contribution from the substrate), which is referenced by suppressing the formation of Cr reinforcement phases, like borides and carbides. In C3 coating, this effect was more pronounced because of the higher dilution coating and consequently higher Fe content. Dilution effect on a Ni-Cr-B-Si coating deposited via laser cladding was also evaluated by Hemmati; Ocelík; De Hosson that also observed a microstructure band with these character istics, describing it as a region of high Fe content and borides absence total [23]. Da Silva; D’Oliveira (2016) and Da Silva; D’Oliveira (2015) evaluated tribological performance at high temperatures of a Ni-Cr-B-Si alloy deposited via PTA. Besides to the described microstructure bands, the authors observed a fourth region (hypereutectic with lower Fe content) [10,24]. In the present research, the failure to obtain a region with this characteristic is due to higher cooling rate imposed by laser cladding, which possibly exceeded that of the PTA, and did not allow sufficient time for its establishment. Another influence source on the microstructure establishment is the temperature gradient/solidification ratio (G/R). According to Kou, near the fusion line, this ratio is high, characterized by a planar solidification front. However, in the solidification course, the ratio is reversed (G re duces and R increases), originating a dendritic growth parallel to fusion line [24,25]. In the literature, several authors describe similar behavior, attributing them mainly to high Fe content and G/R ratio variation [10, 23,24]. Fig. 6 shows central microstructure of the cross-sections, where is possible to note same constituents’ types in all coatings. Difference is concentrated in the distribution, morphology and size of these. Micro structure was established through a dendritic formation, whose distri bution uniformity varied according to dilution, where, compared to C1; C2 and C3 present a matrix with smaller dendritic structures. EDX analysis (Fig. 7) from the microstructure characteristics, indicated: point
� Wear regime defined by the Ha/H ratio (where Ha and H are the abrasive and material hardness, respectively) of Khruschov (1957) [18]. The quartz mineral hardness (1160 HV1) was considered as the abrasive hardness of the sand used. � Volumetric loss: measured by weighting the samples made before and after the tribological tests. From the mass loss values, using Eq. (1), these were converted to volumetric loss. (1)
Q ¼ (ml/ρ) * 1000 3
Where: Q ¼ Volumetric loss (mm ); ml ¼ mass loss (g) and ρ ¼ material density (g=cm3 ). � Evaluation of the dimensional wear coefficient per unit sliding dis tance (k) through the Archard model (1953) [19]: from the volu metric loss measured previously, k coefficient was estimated by the ratio presented in Eq. (2). (2)
k ¼ Q / (LN * D) 3
Where: k ¼ Archard’s wear coefficient [mm /(Nm)]; LN ¼ normal load and D ¼ sliding distance. � Wear surface analysis: SEM technique was used to evaluate the wear tracks formed during the tribological tests. The main purpose of this application was to identify the wear micromechanisms action. 3. Results and discussion 3.1. Microstructure The deposition parameters selection methodology resulted in coat ings with different characteristics. The top surfaces showed uniformity, with absence of voids and other discontinuities. Beads width presented a reduction proportional to the scanning speed (V) increase, even increasing laser power (P). Cross-sections geometric characteristics were very similar to the planned performed. Fig. 3 shows the coatings cross sections visualized by SEM. Fig. 3 reveals regular geometries, without inclusions and cracks. Some lack of metallurgical adhesion points on the C1 coating (high lighted in red) were observed, which did not compromise its integrity and tribological performance at the top surface. This discontinuity _ appearance is related to the energy density. As the powder feed rate (m) was kept constant, the laser beam power (P) employed in C1 was insufficient to melt the entire amount of powder added in the fusion set, together with enough substrate portion, generating a poorer metallur gical adhesion. Coatings C2 and C3 were deposited with higher scanning speed than C1, higher laser beam power, but with a lower energy input, resulting in melting all the added powder in these cases, preventing these discontinuities appearance. Evaluating the surface of both coatings, cracks presence is visible, which vary in size. The cracks arise due to the high tensile residual stresses resulting from the cooling process. Greater P and V resulted in coatings with higher cracks concentration, indicating a higher cooling rate. The result was verified by analyzing the C3 surface (deposited with higher P and V), which revealed a larger quantitative number of cracks, as shown in Fig. 4. In Fig. 4, it is observed that the cracks are oriented perpendicularly to the deposition direction, reinforcing the harmful effect of the residual tensile stresses resulting from the cooling process. Guo et al. evaluated Ni-Cr-B-Si alloy coatings deposited under similar conditions. The au thors observed cracks of this nature appearance and emphasize that the coatings deposited on preheated substrate to 300 � C did not present any �k et al. analyzed cracks [20]. Deschuyteneer et al. (2015) and Vost�ra Ni-Cr-B-Si coatings with and without WC addition as reinforcement 4
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 3. Coatings cross sections obtained via SEM (30x).
analyzed a Ni-Cr-B-Si alloy coating deposited by laser cladding with different V levels (5, 10, 15 and 20 mm/s) and observed that rein forcement phases size reduced gradually for higher V [8]. The results of these experiments are similar to those obtained in this research. Similar to this paper, Deschuyteneer et al. (2015) highlight the presence of a microstructure composed of large Cr agglomerates quan tities dispersed in a Ni-rich matrix. As for the other alloy elements, good distribution uniformity was observed throughout the matrix [7,21]. In the literature, studies that carried out X-ray diffraction (XRD) analyzes on Ni-Cr-B-Si alloys coatings report several types carbides observation. Guo et al. point out the lack of clarity regarding the phases type present in the matrix of this alloys class [20]. The different microstructure characteristics identified directly reflect mechanical properties of hardness and coatings tribological behavior. 3.2. Microhardness The average microhardness values of coatings are shown in Fig. 8. It is interesting to note that, even employing relatively different deposition
Fig. 4. Cooling cracks in the coating C3 surface.
Fig. 5. Different microstructure regions visualized via MO (200x).
1 (Pt1, dark region of Fig. 6) - Cr carbides (5CrC), indicated by arrows; point 2 (Pt2, gray region of Fig. 6) - cut dendritic arm and point 3 (Pt3, light region of Fig. 6) - eutectic matrix composed predominantly of Ni and columnar dendrites, highlighted in red. In the literature, by EDX analysis, Deschuyteneer et al. (2015) report the similar phases obser vation [21]. Table 3 shows chemical composition of the points shown in Fig. 7. In Fig. 6, lower concentration and size of the reinforcement phases (CrC) in the higher dilution coatings is observed. This result is consistent �k et al. and Hemmati; Ocelík; De Hosson, who noticed a with Vost�ra proportional number and size phases reduction with P increase, causing a decrease in microhardness, due to a higher dilution [22,23]. Chen et al.
5
parameters, resulting in significant dilution, microhardness average was similar. However, it is important to highlight that the most extern region of the sample is less affected by dilution, which may justify this simi larity. The more intern regions (in the depth direction) are most affected by the increase in Fe content, which reduces the coatings mechanical resistance. In the literature, there are a number of researches reporting different microhardness results from the laser cladding deposition parameters variation, either individually or together. Chen et al. evaluated different scanning speeds and observed that higher velocity of the analyzed ranges exceeded by more than 200 HV0.2 previous coating microhard ness, a behavior attributed to lower energy amount [8]. This result is compatible with the one observed in this paper since parameterization applied generated similar energy input. Fig. 8 also shows the correlation between microhardness and dilution, where it is possible to observe an
CrC: Cr carbides. 5
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 6. Coatings central region microstructure obtained via MEV (2000x).
inversely proportional behavior. There was a gradual increase from C1 to C2, followed by a larger steeper between C2 and C3. This result is intrinsically linked to its relationship with chemical composition and
coatings microstructure. Regardless of the deposition technique, microstructure exerts influ ence on the microhardness. In the literature, an interesting example is
Fig. 7. EDX of the different microstructure phases identified via MEV. 6
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Table 3 Chemical composition of the different microstructure regions identified. Chemical Element
Point 1 (wt.%)
Point 2 (wt.%)
Point 3 (wt.%)
Carbon Silicon Chromium Iron Nickel
45.0 0.4 45.8 0.9 5.9
43.2 1.8 4.9 3.6 43.8
42.6 1.0 4.4 2.1 47.8
Fig. 10. Microhardness/wear coefficient correlation.
phases as it approaches the substrate. García et al. employed the laser cladding process with deposition parameters and a commercial Ni-Cr-B-Si alloy like this work, with and without WC addition. In the base alloy microhardness profile analysis, analogous to this one, the authors observed higher values in the regions close to the coatings top (about 500 HV1), a result that was also attrib uted to a higher CrC phases concentration in this region [26]. Kaiming et al. evaluated the Ni45 alloy resistance, which chemical composition is similar to this work. Deposition parameters used by the authors resulted in a higher microhardness in the HAZ (600 HV0.5), while in the coating (500 HV0.5). This result has been attributed to the occurrence of quenching in this region during coating deposition [27]. Observing at the average surface microhardness, the difference be tween coatings is relatively small. This result is probably related to similar microstructure characteristics in this region. As the same phases’ types were identified and the microindentation hardness test was quite punctual, no significant difference was identified. This result also asserts the high integrity of coatings. Higher variation was observed in the microhardness profiles, where elevated thermal gradient processed coatings presented inferior performance, due to the higher HAZ, greater planar microstructure band (with a higher Fe content and lower hard ness) and lower concentration and size of reinforcement phases dispersed in its matrix.
Fig. 8. Microhardness/dilution relation.
3.3. Wear results
Fig. 9. Microhardness profiles.
3.3.1. Volumetric loss Table 4 shows abrasive wear regime types identified through the Khruschov Ha/H ratio (1957) [18]. In the coatings tested via dry san d/rubber wheel test, the ratio presented a value higher than the limit of 1.2. This result indicates that all coatings were subjected to a severe wear regime. As the difference between the surface microhardness average was small, the ratio presented close values. With the results of volumetric loss that is proportional to the k wear coefficient, it was observed:
presented by Da Silva; D’Oliveira (2015), who compared microstructure and microhardness of Ni-Cr-B-Si coatings deposited by PTA and laser cladding. The authors concluded that in both techniques, the dilution significantly impacted the microstructure and coatings microhardness. The different cooling rate between the two techniques directly modified the microstructure [24]. Fig. 9 shows average microhardness profiles. The coatings presented an average higher to 550 HV1, value almost four times higher to sub strate (150 HV1). A considerable decrease is observed near the sub strate/coating interface, including diluted region and HAZ, respectively. Microhardness profiles followed microstructure variation, where three levels were verified, corresponding to the three microstructure regions. Microhardness increased proportionally to the CrC concentration elevation towards coatings top. In contrast, a proportional decrease of this was observed in relation to reduction in the amount and size of CrC
� In the first 5 min of test, the wear is high, which is probably attrib uted to the following factors: contact area increases between the specimens; vibrations at the moment of wear interface establishment and oxide layer remove. This last phenomenon was an issue explored by Bressan; Hesse; Silva Jr. (2001), who highlight the same as being one of the main oscillations causes in the initial tribological tests phase [28]. Besides to the mentioned characteristics, coatings wear had an additional influence source, which are the cooling cracks present on the sample’s surfaces. During the tests, these acted as anchoring points, generating coating fragments detachments (sub ject to be discussed in more detail in topic 3.3.3).
Table 4 Wear regime by Khruschov’s ratio. Coating
C1
C2
C3
Ha/H ratio Regime type
2.0 Severe
2.1 Severe
2.2 Severe
7
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 11. Features worn out tracks.
� After the stabilizing step, there is a drop in the volumetric loss for 10 min, indicating the wear transition. This occurs due that valleys left by removed fragments are being filled by abrasive material that is not effectively removing coating material. � A new material detachment occurs close to 15 min of the test, evi denced by another high volumetric loss point. � For 20 min of test, a behavior similar to 10 min was observed, probably related to the same reason. � Between 25 and 30 min, average loss and standard deviation differ ence are small compared to the other pauses performed in the test, indicating the uniform volumetric loss regime establishment. The average accumulated volumetric loss over the 30 min of test showed an inversely proportional correlation to microhardness. C1 showed the lowest volumetric loss, followed by C2 and C3, respectively. The matrix characteristics and reinforcement phases (CrC) were the most significant influence sources. The higher concentration and phase size of CrC showed a positive effect on the microhardness and conse quent resistance to abrasive wear. The correlation analysis of hardness and volumetric loss shows a better performance of the coatings with the highest concentration and size of CrC phases. This result is because these coatings are less affected by the sand action when compared to the matrix, which has a lower hardness. Due to the larger dimensions, these phases tend to stay fixed longer in the matrix, giving it a longer protection time. These charac teristics justify the better performance of C1. In the literature, Deschuyteneer et al. evaluated the abrasion resis tance of a Ni-Cr-B-Si alloy with and without WC addition by the dry €kki; Vouristo (2009) [29], sand/rubber wheel test [7]. Nurminen; Na they analyzed wear resistance of various Metal Matrix Composites (MMC) coatings deposited via laser cladding, through the dry san d/rubber wheel test. Alloy studied by the first author and one of the alloys evaluated by the second author is similar to this work. Both au thors highlight obtention of higher wear resistance in coatings that presented higher concentration and size of reinforcement phases. 3.3.2. Archard wear coefficient Fig. 10 shows an inversely proportional correlation between micro hardness and k wear coefficient. The results also expose the stability of this coefficient. Throughout the test, the wear depth increases and other layers of microstructure are affected. As the k coefficient average remained constant, the microstructure variation along depth was not enough to generate significant changes in the coatings. However, the different microstructure characteristics between the three coating con ditions influenced in the k wear coefficient. Moreover, at C2 and C3, wear level was sufficient to reach the hypoeutetic bands and even the planar region, which have lower resistance in relation to the eutectic range. In
Fig. 12. Wear micromechanisms identified via MEV (1500x).
8
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 13. Wear micromechanisms interruption (1500x).
this case, it is noticed that C3 indicate of inability to withstand more severe test conditions since their inverse correlation with microhardness presents practically nonexistent. 3.4. Wear surfaces 3.4.1. Wear micromechanisms In general, volumetric loss and the respective k wear coefficients were more pronounced in the coatings with higher dilution, higher cracks concentration, lower amount and size of CrC phases and conse quent lower microhardness. These results reflected directly on the worn surfaces aspects. Abrasive wear mechanism action was observed through grooves (Fig. 11) formed on the sample’s surfaces. An interesting point is that these grooves presented deeper in the regions between beads, a result that can be related to the greater Fe migration to the top of the coatings through the convection currents resulting from solidification process. The fact of this effect is more pronounced in the higher dilution coatings (higher Fe content) reinforces this justification, indicating that these are also lower microhardness regions. During the tests, sand flow tend re mains trapped at these regions, increasing wear. In the literature, Deschuyteneer et al. report similar behavior in the worn surfaces ana lyzes of their coatings [7]. In this research, wear micromechanisms types identified were com mon to all coatings, and the difference was concentrated in their quantitative and severity action. C1 surface presented less severe wear aspects when compared to the other coatings, a result related mainly to its higher microhardness, which prevents the penetration of abrasives on the surface, reducing material removal. Abrasive wear predominant action was observed through microploughings and microcuttings. There are also microindentations characteristic of the wear test and CrC phases fractured by the abrasive particles action. In the specimen’s central re gion, all coatings presented the microploughing as predominant wear micromechanism. The microcutting action was inhibited in regions with CrC phases presence, a more prominent aspect was observed in C1. Fig. 12 shows these worn surfaces. In some points, transitions from microploughing to microcutting are observed, going back to the wear micromechanisms transition model of Zum Gahr [30]. By comparing results of Fig. 12 with k wear coefficient, worn surfaces aspects were consistent, being more severe in C2 and C3, which pre sented performance inferior to C1. However, the intensity degree of micromechanisms identified action was relatively contained. In this way, a more detailed analysis was carried out on the influence of CrC phases and cooling cracks present in the coatings.
Fig. 14. Worn surfaces close to the cooling cracks obtained via MEV: C2 (500x) and C3 (1000x).
3.4.2. CrC carbides and wear surfaces The coatings surfaces presented wear tracks arranged in random directions. No CrC phases detachments of the matrix were observed, and its protective action was identified in several regions, more highlighted on the C1 surface. In Fig. 13, regions where CrC phases act as barriers, preventing the abrasive particles advance on the coating surface, are clearly observed. Besides that, there are points where CrC phases could not avoid, however, they made difficult the continued abrasive sand action flow, resulting in the less severe microploughings formation. In the literature, most of the authors that used dry sand/rubber wheel test highlight an increase in the coatings abrasion resistance through carbides concentration increasing [7,13,21,29,31,32]. Similar to behavior observed in this work, the authors highlight as advantages hardness enhancement, which increases the surface resistance to abra sive particles penetration, the restraint effect that these imposes to the wear micromechanisms advancement, besides the reduction of matrix area in direct contact with wear interface. However, increase in the reinforcing phases concentration is only interesting up to a certain limit. In this sense, Chatterjee; Pal evaluated a ratio of wear rate/carbide concentration and describe that from a given point, the excess carbides became harmful, where the number of frac tured phases acting as abrasive particles increased too much [33].
9
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
3.4.3. Cooling cracks and wear surfaces Coatings with higher cooling cracks concentration showed higher wear. SEM analyses (Fig. 14) revealed coating material detachments in the cracks and regions close to these ones. This behavior was more evident on the C3 surface when compared to C2. On the other hand, C1 did not present this problem, a result that can be attributed to the lower amount and intensity of cooling cracks identified on its surface. The more highlighted cracks concentration in C2 and C3 is due to the com bination between higher P and V used in these that resulted in a higher thermal gradient. In Fig. 14C2, a crack is observed parallel to the deposition and dry sand/rubber wheel test direction, respectively, due to the tensile stresses resulting from the cooling process. It is possible to observe coating material detachment in the cracks and in nearby regions. Besides that, above the removed region, the eminence of a new coating material detachment is visible, which would increase the wear tracks formed and respective side effects of this phenomenon. This behavior tends to occur cyclically as the test proceeds, generating larger and tending to a cascade effect. Fig. 14C3 presents same behavior, however, in a more severe and deconcentrated way, where along the length of the entire crack, there are detachments, making it even difficult to identify regions that only show cracks. Considering the Fig. 14C2 and Fig. 14C3, it is possible to infer that wear identified on C2 surface is more specific and severe than on C3. However, in C2, cracks collapsed in a restricted region, while in the second, effect covered much of the surface. Higher dilution and conse quent lower microhardness of C3, together with its higher cracks con centration, are consistent with this result. In the literature, works that employ most varied deposition process types, report similar problems. For example, Colaço; Maranho who employed the flux-cored arc welding coating process, as also Natarajan et al. who used laser cladding tech nique report that microploughings formed by the abrasive sand sliding removed large coating material fragments in the regions near cooling cracks, increasing wear [13,32]. Although of being different cases with a similar purpose, the key problem indicated is basically the same, which also extends to the respective work. In theory, material detachment close to the cracks can be charac terized as brittle rupture. However, in this work, it was not possible to confirm microcracking action associated with this behavior, since these cracks origin is known. The large coating fragments loss, caused by material breakage at the cracking limit is also indicated as a key factor in the work of Houdkov� a et al. [12]. Another point observed on both coatings surface is some dark stains, from the abrasive sand accumulation, where a portion of these was fractured and embedded in the surfaces of the coating, acting as a pro tective film, contributing to wear reduction. These stains origin was confirmed by EDX carried out in the region, which indicated a high content of Si (20%) and oxygen O (14%), silica constituent elements (SiO2). Coronado Marin observed similar stains on the AISI 1045 steel surface after performing dry sand/rubber wheel test, attributing them to the same justification [31]. Evaluating the general worn surfaces aspects, the positive influence of the CrC phases presence and the highly cooling cracks damaging ef fects, it is possible to infer that the microploughing action associated to the brittle material detachment parallel to the cooling cracks were predominant causes of the observed wear levels. Besides that, micro cuttings and microindentations presence can also be indicated as responsible for abrasive wear rate.
�
�
�
�
integrity and tribological performance. According to the powder feed rate, parameters used to this coating gave it a lower dilution in comparison to the others. Regions with highest cooling cracks observed in C2 and C3 were attributed to the higher thermal gradient and these perpendicular arrangement defects confirmed the damaging effect of the tensile stresses resulting from cooling process. The microstructure formed by the chemical composition gradient and G/R ratio was constituted by CrC phases dispersed in a Ni den dritic eutectic matrix, whose size and distribution varied according to the dilution. With the surface and cross sections analyzes of C1, was identified a higher microhardness average, that is due to modifying mechanical and metallurgical properties. Microhardness profile showed three different levels, corresponding to the three different microstructure regions. Establishing a magnitude order from higher to lower abrasive wear resistance, coatings can be classified: C1, C2 e C3, a result that can be attributed mainly to microstructure and microhardness characteris tics, which presented inversely proportional with the results of volumetric loss and k wear coefficient. In the worn surfaces analyzes, a considerable wear micromechanisms part were interrupted for the restraint effect formed by CrC phases. It is possible to infer that the regions with higher material volume removed, was due to cooling cracks. In general, it was observed a better tribological performance of C1 in comparison to the other coatings, within the analyzed context, highlighting it to the most suitable for more severe wear conditions applications. However, depending on the solicitation degree, C2 and C3 are also reproducible, since the higher concentration and severity of the cooling cracks identified in these can be reduced through substrate preheating treatments.
Declaration of competing interest None. Acknowledgements The Federal University of Santa Catarina (UFSC, BRA) and SATC College (SATC, BRA) by scientific technical support. The Coordination for the Improvement of Higher Education Personnel (CAPES, BRA) and National Council for Scientific and Technological Development (CNPq, BRA) for financial support. References [1] H€ arkisaari P. Wear and friction effects on energy consumption in the mining industry. Master of science thesis. Tampere: Tampere University of Technology; 2015. [2] Holmberg K, Erdemir A. Influence of tribology on global energy consumption, costs and emissions. Friction 2017;5:263–84. https://doi.org/10.1007/s40544-0170183-5. [3] Holmberg K, Kivikyt€ o-Reponen P, H€ arkisaari P, Valtonen K, Erdemir A. Global energy consumption due to friction and wear in the mining industry. Tribol Int 2017;115:116–39. https://doi.org/10.1016/j.triboint.2017.05.010. [4] Toyserkani E, Corbin S, Khajepour A. Laser cladding. Boca Raton (USA): CRC Press LLC; 2005. [5] Zhong M, Liu W. Laser surface cladding: the state of the art and challenges. J Mech Eng Sci 2010;224:1041–60. https://doi.org/10.1243/09544062JMES1782. [6] Davim JP. Laser in manufacturing. Hoboken: John Wiley & Sons; 2013. [7] Deschuyteneer D, Petit F, Gonon M, Cambier F. Influence of large particle size–up to 1.2 mm–and morphology on wear resistance in NiCrBSi/WC laser cladded composite coatings. Surf Coat Technol 2017;311:365–73. https://doi.org/ 10.1016/j.surfcoat.2016.12.110. [8] Chen J, Li J, Song R, Bai LL, Shao JZ, Qu CC. Effect of the scanning speed on microstructural evolution and wear behaviors of laser cladding NiCrBSi composite coatings. Opt Laser Technol 2015;72:86–99. https://doi.org/10.1016/j. optlastec.2015.03.015. [9] Luo X, LI J, LI GJ. Effect of NiCrBSi content on microstructural evolution, cracking susceptibility and wear behaviors of laser cladding WC/Ni–NiCrBSi composite coatings. J Alloy Comp 2015;626:102–11. https://doi.org/10.1016/j. jallcom.2014.11.161.
4. Conclusions From the results obtained, it was possible to reach some conclusions: � The coatings presented a good surface finish for shape quality. Points with lack of metallurgical adhesion identified in C1, attributed to the insufficient energy density parameter did not compromise its 10
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002 � anek C. The influence of process �k M, Houdkov� [22] Vost�ra a S, Bystrianský M, Ces� parameters on structure and abrasive wear resistance of laser clad WC-NiCrBSi coatings. Mater Res Express 2018;5. https://doi.org/10.1088/2053-1591/aad859. [23] Hemmati I, Ocelík V, De Hosson JTM. Dilution effects in laser cladding of Ni–Cr–B–Si–C hardfacing alloys. Mater Lett 2012;84:69–72. https://doi.org/ 10.1016/j.matlet.2012.06.054. [24] Da Silva LJ, D’Oliveira ASCM. Liga NiCrSiBC: microestrutura e Dureza de Revestimentos Processados a Arco e a Laser. Soldagem & Inspeç~ ao. 2015;20:39–47. https://doi.org/10.1590/0104-9224/SI2001.05. [25] Kou S. Welding metallurgy. New Jersey: John Wiley & Sons; 2003. [26] García A, Fern� andez MR, Cuetos JM, Gonz� alez R, Ortiz A, Cadenas M. Study of the sliding wear and friction behavior of WC þ NiCrBSi laser cladding coatings as a function of actual concentration of WC reinforcement particles in ball-on-disk test. Tribol Lett 2016;63:41. https://doi.org/10.1007/s11249-016-0734-3. [27] Kaiming W, Yulong L, Hanguang F, Yongping L, Zhenqing S, Pengfei M. A study of laser cladding NiCrBSi/Mo composite coatings. Surf Eng 2016;34:267–75. https:// doi.org/10.1080/02670844.2016.1259096. [28] Bressan J, Hesse R, Silva Jr E. Abrasive wear behavior of high speed steel and hard metal coated with TiAlN and TiCN. Wear 2001;250:561–8. https://doi.org/ 10.1016/S0043-1648(01)00638-X. [29] Nurminen J, N€ akki J, Vuoristo P. Microstructure and properties of hard and wear resistant MMC coatings deposited by laser cladding. Int J Refract Metals Hard Mater 2009;27:472–8. https://doi.org/10.1016/j.ijrmhm.2008.10.008. [30] Zum Gahr K-H. Microstructure and wear of materials, tribology series. New York: Elsevier Science; 1987. [31] Coronado Marin JJ. Efeito do tamanho do abrasivo no desgaste de metais. Tese de Doutorado. S~ ao Paulo: Universidade de S~ ao Paulo; 2010. [32] Colaço FHG, Maranho O. Avaliaç~ ao da perda de massa de revestimento duro depositado por soldagem com arame tubular de liga FeCrC-Ti. Soldagem & Inspeç~ ao. 2014;19:58–68. https://doi.org/10.1590/S0104-92242014000100008. [33] Chatterjee S, Pal T. Wear behaviour of hardfacing deposits on cast iron. Wear 2003; 255:417–25. https://doi.org/10.1016/S0043-1648(03)00101-7.
[10] Da Silva LJ, D’OLIVEIRA ASCM. NiCrSiBC coatings: effect of dilution on microstructure and high temperature tribological behavior. Wear 2016;350: 130–40. https://doi.org/10.1016/j.wear.2016.01.015. [11] Miguel J, Guilemany J, Vizcaino S. Tribological study of NiCrBSi coating obtained by different processes. Tribol Int 2003;36:181–7. https://doi.org/10.1016/S0301679X(02)00144-5. � Smazalov� �k M, Schubert J. Properties of NiCrBSi coating, as [12] Houdkov� a S, a E, Vost�ra sprayed and remelted by different technologies. Surf Coat Technol 2014;253: 14–26. https://doi.org/10.1016/j.surfcoat.2014.05.009. [13] Natarajan S, Anand EE, Akhilesh KS, Rajagopal A, Nambiar PP. Effect of graphite addition on the microstructure, hardness and abrasive wear behavior of plasma sprayed NiCrBSi coatings. Mater Chem Phys 2016;175:100–6. https://doi.org/ 10.1016/j.matchemphys.2016.02.076. [14] Yin B, Zhou HD, Chen JM, Yan FY. Microstructure and high temperature sliding wear behaviour of laser remelted NiCrBSi/35(WC-12Ni) composite coating. Surf Eng 2011;27:458–63. https://doi.org/10.1179/026708410X12687356948517. [15] Ocelík V, Nenadl O, Palavra A, De Hosson JThM. On the geometry of coating layers formed by overlap. Surf Coat Technol 2014;242:54–61. https://doi.org/10.1016/j. surfcoat.2014.01.018. [16] ASTM-E92. American society for testing and materials: standard test methods for Vickers hardness and knoop hardness of metallic materials. 2017. [17] ASTM-G65. American society for testing and materials: standard test method for measuring abrasion using the dry sand/rubber wheel apparatus. 2016. [18] Khruschov MM. Resistance of metals to wear by abrasion, as related to hardness. Conf. Lubrication and Wear 1957:655. [19] Archard J. Contact and rubbing of flat surfaces. J Appl Phys 1953;24:981–8. [20] Guo C, Zhou J, Chen J, Zhao J, Yu Y, Zhou H. High temperature wear resistance of laser cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings. Wear 2011;270: 492–8. https://doi.org/10.1016/j.wear.2011.01.003. [21] Deschuyteneer D, Petit F, Gonon M, Cambier F. Processing and Characterization of laser clad NiCrBSi/WC composite coatings - influence of microstructure on hardness and wear. Surf Coat Technol 2015;283:162–71. https://doi.org/10.1016/ j.surfcoat.2015.10.055.
11
Contents lists available at ScienceDirect
Tribology International journal homepage: http://www.elsevier.com/locate/triboint
Abrasion resistance of Ni-Cr-B-Si coating deposited by laser cladding process☆ � de Sousa a, *, Francisco Ratusznei a, Milton Pereira a, Jurandir Marcos Sa Richard de Medeiros Castro b, Elvys Isaías Mercado Curi b a b
Laborat� orio de Mec^ anica de Precis~ ao, Universidade Federal de Santa Catarina, Florian� opolis, Brazil Laborat� orio de Vibraç~ oes e Tribologia, Faculdade SATC, Criciúma, Brazil
A R T I C L E I N F O
A B S T R A C T
Keywords: Abrasion resistance Chromium carbides Thermal gradient Cooling cracks
Ni-Cr-B-Si alloy deposited by laser cladding has its tribological performance evaluated. Deposition parameters selection is a difficult task. In this work, Ni-Cr-B-Si coatings were deposited by a fiber laser source with a coaxial powder nozzle on low carbon steel substrate. Microstructure, microhardness and abrasive wear tests were per formed. Results showed coatings with good surface adhesion. Microhardness was 10% higher in coatings with higher concentration and phase size of chromium carbides. Volumetric loss and wear coefficient showed 10% variation between the coatings. Coatings with different thermal gradients showed different dilution levels, affecting their abrasion resistance. The better tribological performance was obtained for coatings with lower cooling cracks density and higher volumetric fraction of carbides, which both mitigated wear micromechanism action.
1. Introduction Currently, surface engineering has proposed new technologies and alloys to enhance the performance of mechanical components. The materials subjected to severe wear require coatings that increase the lifetime of machine elements, and consequently, minimize financial losses due to shutdowns for the maintenance and replacement of worn components. Among the various types, abrasive wear mechanisms are the most frequently encountered, resulting from the friction and sliding between the hard particles and the surface of solid components. Industry accounts for approximately 30% of total world energy consumption [1]. About 23% of these 30% come from tribological contacts [2]. Sectors of great economic importance, such as the agri cultural, oil and gas and mining industry, requires and are interested in increasing their components lifetime. In mining, for example, 40% of the expenses related to friction losses, 27% of the deteriorated parts replacement, 26% of the workforce and 7% of production losses can be avoided using efficient tribological solutions [3]. The environment can also be greatly benefited by advanced and efficient tribological solutions implementation. An extensive study has shown that in the short term (approximately 8 years), on a global scale,
due to the implementation of advanced tribological technologies, CO2 emission reduction can be as much as 1.460 Mt, representing a saving of 455.000 € (Euros). In the long term (around 15 years), these values rise to 3.140 Mt and 970.000 €, respectively [2]. According to this context, it is important to improve the tribological properties, especially abrasive wear resistance, which causes volume loss, dimensional changes and failures in mechanical and related com ponents. Thus, the search for alternative manufacturing processes that provide cost-effective features has become a constant aim for research centers around the world. One adopted solution is the application of hard coatings by deposi tion processes, such as thermal spraying and arc welding processes. In these techniques, layers with special properties are deposited over a lower mechanical strength substrate. Instead of manufacturing parts composed entirely by special materials, only the surface area of interest is coated, and this approach can drastically reduce the components final cost. Despite the cited processes can generate acceptable results, there are still challenges to be overcome to allow high quality coatings production without problems related to distortions appearance, lack of metallur gical adhesion between substrate and addition material, extensive heat
This Work was presented at TRIBOBR2018. * Corresponding author. E-mail address: [email protected] (J.M.S. Sousa).
☆
https://doi.org/10.1016/j.triboint.2019.106002 Received 4 August 2019; Received in revised form 1 October 2019; Accepted 3 October 2019 Available online 16 October 2019 0301-679X/© 2019 Elsevier Ltd. All rights reserved.
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
affected zone (HAZ) and high dilution level [4–7]. In order to solve these problems commonly related to conventional arc and spraying processes, several studies have been carried out involving other solutions. Laser cladding technique, for instance, has the ability to overcome these challenges, generating better coatings when compared to other deposition methods. A number of these features is provided by the unique properties from laser radiation, which are: high directionality beam, high power available, high versatility and achiev able energy density, which allow the processing of a wide variety of materials and alloys, including the processing of complex geometry components [4,6]. Nickel (Ni) based alloys are characterized by having good abrasion and corrosion resistance properties, even in high-temperature environ ments. Among these, Ni-Cr-B-Si class alloys, which have good resistance to the mentioned wear mechanisms, also presents high hardness values [8–10]. These alloys are classified as self-fluxing, a term referring to alloys capable of reducing their liquidus temperature, allowing the hard phases precipitation. Each constituent element of these alloys has a specific purpose: Cr (chromium) protects against corrosion at high temperatures and inserts hard phases; B (boron) reduces the melting point and also contributes to the hard phases increase; Si (silicon) is responsible for self-fluxing properties and C (carbon) inserts wear-resistant carbides [11–13]. Ni-based alloys are widely applied in components from various in dustries, such as the oil and gas industry, glass molds, cold metalwork components and slurry purge elements. Besides to the cited examples, Ni-Cr-B-Si class alloys are highly suitable for maintenance and repair in industrial environments, such as extrusion screws for the plastic in dustry [13,14]. Such alloys are commonly deposited by thermal spraying processes. However, besides to the previously mentioned problems for these techniques, coating layers with high porosity concentration are also generated [11]. To fill this gap, laser cladding process is being applied, which provides the desired characteristic through a highly concentrated beam, mitigating the porosity while inserts less heat on the substrate. Laser cladding characteristics, combined with the choice of materials with good mechanical properties and suitable processing parameters, allow the high-quality coatings production. However, parametrizing the process adequately is a complex task, since there are an extensive number of variables involved, making the process multidisciplinary, requiring knowledge in the physics and metallurgy of the deposition process and also materials science. This paper aims to contribute to this discussion by evaluating the abrasion resistance of Ni-Cr-B-Si alloy, deposited by laser cladding with different processing parameters levels on AISI 1020 carbon steel sub strate, applying the wear abrasion test by using the standard ASTM G65. The processed coatings will be subjected to dilution, microhardness, microstructure, volumetric loss, Archard wear coefficient and worn surfaces aspects analysis. The main objectives are to evaluate the deposition variables effects, the characteristics and properties of the alloy on the coatings obtained quality, besides establishing a processing methodology for the manufacture of coatings with this deposition technique and alloy.
(76.0 � 25.4 � 10.0 mm) were prepared, with average roughness Ra ¼ 0.7 μm (test surface). Ni-Cr-B-Si alloy 1545-00 (powder) from €gana €s S.A manufacturer was used as the coating material. The alloy is Ho suitable for coatings, has a hardness value of 490 HV30 and a density of 7.1 g/cm3. The particle powder size of 150.0 þ 53.0 μm, with rounded morphology, was observed by scanning electron microscopy (SEM). Due to the laser cladding nozzle and powder feeder requirements, alloy was sifted to the granulometric range of 106 þ 53.0 μm. Table 1 shows its chemical composition. 2.2. Coatings deposition process Prior to the deposition process, substrate samples were subjected to an abrasive blasting process (CMV GS-9075X equipment with SINTER BALL NORBLAST sand) in order to reach an opaque surface, avoiding reflectivity, to reach a uniform material absorptivity during processing. Thereafter, they were cleaned in ethyl alcohol to ensure impurities removal. Samples were coated using an Yb-Ytterbium fiber laser source from the IPG PHOTONICS® manufacturer, with maximum power of 10-kW, wavelength (λ) ¼ between 1070 nm and 1080 nm, laser beam diameter at beam waist ¼ 0.85 mm, M2 ¼ 1.1. Specimens are moved using a numerically controlled XYZ coordinate table. Coating material is added by a disk type powder feeding system (GTV) with a1powder feed rate _ ¼ between 0.5 and 30.0 g/min. Powder feeder nozzle is the coaxial (m) type. Argon (Ar) was used as protection and carrier gas. Deposition parameters were determined based on 24 single beads preliminary tests, where2laser beam power (P) and3scanning speed (V) were varied, keeping the other parameters constant. Fig. 1 flowchart simplifies the steps of this approach. Coated specimens resulting from this step were subjected to a standard metallography procedure: sanding (80, 220, 400, 600 and 1200 Mesh), polishing with 0.3 μm alumina and chemical etching in 2.0% nital reagent for 3.0 s. These specimens were subjected to optical microscopy. Using the Image J software, single beads morphology was evaluated and the ones with lowest dilution values, without defects and/or discontinuities such cracks, pores, de tachments at the interface substrate-coating, and with better geometric aspects, were selected. The methodology considered: height (h � 0.3 mm), width (w � 1.0 mm), area, substrate penetration (p lowest possible) and wettability angle (β < less than 90� to avoid pore insertion in the overlap step). The final deposition parameter was the overlapping condition. The overlap rate was estimated from predictions provided by a code pro grammed into MATLAB software, which was developed based on the work presented by Ocelík et al. [15]. After inserting the input data (modeling parameters: clad width - mm, clad height - mm, clad overlap %, number of clad tracks and layers) into the algorithm and performing the interactions, cross sections with the best geometric surface and resulted coating height were obtained. In this way, the final deposition parameters selection was based on the relation between the following aspects: 1 - lower dilution values; 2 no cracks; 3 - absence or fewer pores presence; 4 - geometric Table 1 Ni-Cr-B-Si 1545-00 alloy chemical composition.
2. Methodology
Elements (wt.%)
2.1. Substrate and coating material AISI 1020 laminated steel with chemical composition wt.%: 0.18–0.23 (C); 0.3–0.6 (Mn); 0.04 (P); 0.05 (S) and (Fe) in balance was used as the substrate material. The choice of this material is due to its wide industrial applicability. In this way, its choice adequately simulates a real application condition. The substrate was subjected to machining process (cutting, milling and grinding) in order to satisfy the dimensions and surface quality required by ASTM G65 standard. Rectangular specimen
C
Fe
Cr
B
Si
Ni
0.3–0.4
2.0–3.4
7.5–10.0
1.7–2.0
3.3–3.9
Bal.
1 2 3
2
_ powder feed rate. m: P: laser beam power. V: scanning speed.
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 1. Deposition parameters determination.
characteristics and 5 - overlap rate prediction. 3 of the 24 parameter combinations were chosen to deposit the coatings. These are presented in Table 2, being4coatings 1 (C1), 2 (C2) and 3 (C3), respectively. These parameters shown in Table 2 were combined with fixed pa _ ¼ 10.4 g/min); laser beam diameter rameters, such: powder feed rate (m at beam waist (0.85 mm), distance from substrate surface to beam waist (35.0 mm), shielding gas flow (Ar ¼ 15.0 l/min) and carrier gas flow (Ar ¼ 5.0 l/min), to deposit the coatings. For each test group, two sam ples with rectangular area (20 � 25 mm) were deposited to be tested and analyzed (dilution, microhardness and microstructure).
coating), cross-sectional samples were evaluated by optical microscopy (OM) and scanning electron microscopy (SEM) using Backscattering Electron Detector (BSE) a HITACHI device, model TM-3030. Specific regions (presented particular or abnormal aspects) had their chemical composition evaluated by X-ray dispersive energy spectroscopy (EDX) in the SWIFTED 3000 apparatus, coupled to the SEM. 2.4. Tribological test For the tribological tests, two samples of each coating condition were prepared. In order to qualify the Ni-Cr-B-Si 1545-00 alloy abrasion resistance deposited by laser cladding, the dry sand/rubber wheel abrasion test was carried out, as stated by ASTM G65 (2016) [17]. The specimens were cleaned using ethyl alcohol and weighted in a precision scale (MARTE AD500 of resolution 0.001 g) before the start and after each test section in order to measure the volumetric loss due to tribological tests. The tribometer was previously calibrated according to the ASTM G65 requirements for sliding speed, operating stability and abrasive mass flow. The sand used as abrasive underwent a sieving process to reach a 30 Mesh grain size, followed by heating at 100 � C for 1 h in order to remove the moisture. Using ASTM G65 procedure A, the following parameters were selected: 130 N load, 200 rpm speed, 30 min test time, totaling a sliding distance of 4309 m, under a 390 g/min continuous sand flow. In order to perform an analysis of wear tendency with respect to time, the following procedure was performed: each sample was continuously tested for 5 min; after every 5 min of testing, the sample was cleaned again and weighed; after this step, the sample was reposi tioned in the tribometer. This process was performed six times for each specimen. Fig. 2 summarizes the main procedures carried out in the samples: (a) coating as deposited condition, (b) as grinded condition and (c) worn surface after tribological test.
2.3. Mechanical and microstructural characterization Mechanical characterization was carried out with the selection of one sample from each coating condition, randomly chosen. Dilution was measured in the final coatings cross-sections in a procedure analogous to performed on the single beads’ specimens. Coatings hardness was evaluated through Vickers microhardness test, following the specifications indicated by ASTM E92 (2017) [16]: parameters regarding strength, test time and a minimum distance be tween indentations, using the SHIMADZU HMV-2TADW equipment. The HV1 scale with a 10.0 N load for 10.0 s, was used. For the hardness evaluation of the coatings surface, an indentation pattern was elabo rated (in order to obtaining a reliable and low standard deviation result): 20 uniformly distributed indentations (5 mm spacing between each indentation) were made in a rectangular pattern (5 � 4) along the sample length, in order to sweep the region subsequently submitted to the tribological test. To evaluate the hardness profile, following pattern was adopted: 3 indentation lines spaced in 6 mm each other were drawn in the sample width direction (total sample width: 25 mm). The average coatings thickness was 1.5 mm, so 5 indentations were spaced 0.4 mm across the sample height, so that it was possible to reach the different areas (coating, diluted region, HAZ and substrate). Preparation process for metallographic analysis was similar to that used in the single beads preliminary tests. The only difference was the etching, which was performed by using a reagent composed of 10 ml FeCl3, 20 ml HCl, 50 ml H2O and 1 ml HNO3 for 90 s. In order to evaluate the different microstructure areas (substrate, diluted region, HAZ and
2.5. Wear analysis Results from the ASTM G65 tribological tests were evaluated based
Table 2 Coatings deposition parameters. Parameters
C1
C2
C3
Laser beam power (kW) Scanning speed (mm/s) Overlap (%) Energy imput (J/mm) Dilution (%)
1.05 5.0 30 210.0 4.1
1.40 21.7 30 64.6 12.6
1.75 30.0 30 58.3 27.4
4
Fig. 2. Main steps applied to samples: (a) as coated, (b) as grinded and (c) after tribological test.
Coating 1 (C1), coating 2 (C2) and coating 3 (C3). 3
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
on the following perspectives:
phase and also identified cooling cracks, which ceased after preheating the substrate at 400 � C and 350 � C, respectively [21,22]. Besides to the cited researches, there are other authors in the literature that highlight similar results. It is important to note that cracks can also be nucleated from defects in microstructure or inclusions in the material. In the present research, as no impurities of this nature were detected, observed cracks were conferred to the cooling phenomena. Coatings dilution was also consistent with deposition parameters employed and the single beads preliminary tests: C1 ¼ 2.6%, C2 ¼ 4.8% and C3 ¼ 22.6%. It is complex to infer about the deposition/dilution _ parameters relationship. It is likely that the result is connected to m used. As it was maintained constant, it is believed that the higher V used in C2 and C3 provided lower powder insertion per time unit and, even increasing P, an optimum condition was not reached, where the action of both three parameters generated coatings with similar geometric and metallurgical aspects. The dilution impacted the chemical composition, a factor evidenced by EDX analysis, which revealed an increase in Fe content and Ni reduction in the higher dilution coatings. In relation to Cr, B and Si, differences were not significant. In the literature, several papers describe hardness and wear resistance reduction proportional to the coating’s dilution increase, indicating the substrate chemical composition and deposition technique as key factors. Several authors emphasize that this behavior is the major responsible for hardness decreasing, a factor consistent with this research results [8,10,23]. Besides to the proper metallurgical adhesion, three microstructure characteristic regions can be visualized: a precipitate free band (planar), located near the fusion line, a hypoeutectic and a eutectic layer. These changes resulted from chemical composition variation, especially Fe content, where, in the planar and hypoeutectic (high concentration) and eutectic regions (lower content towards the coating top). Fig. 5 (ob tained via OM) indicates these different characteristic regions. In Fig. 5, dilution effect was evident through the precipitate free band establishment, which is larger (has a greater thickness) in C3. This characteristic is possible connected to variation in chemical composi tion, more specifically, the Fe content (the main contribution from the substrate), which is referenced by suppressing the formation of Cr reinforcement phases, like borides and carbides. In C3 coating, this effect was more pronounced because of the higher dilution coating and consequently higher Fe content. Dilution effect on a Ni-Cr-B-Si coating deposited via laser cladding was also evaluated by Hemmati; Ocelík; De Hosson that also observed a microstructure band with these character istics, describing it as a region of high Fe content and borides absence total [23]. Da Silva; D’Oliveira (2016) and Da Silva; D’Oliveira (2015) evaluated tribological performance at high temperatures of a Ni-Cr-B-Si alloy deposited via PTA. Besides to the described microstructure bands, the authors observed a fourth region (hypereutectic with lower Fe content) [10,24]. In the present research, the failure to obtain a region with this characteristic is due to higher cooling rate imposed by laser cladding, which possibly exceeded that of the PTA, and did not allow sufficient time for its establishment. Another influence source on the microstructure establishment is the temperature gradient/solidification ratio (G/R). According to Kou, near the fusion line, this ratio is high, characterized by a planar solidification front. However, in the solidification course, the ratio is reversed (G re duces and R increases), originating a dendritic growth parallel to fusion line [24,25]. In the literature, several authors describe similar behavior, attributing them mainly to high Fe content and G/R ratio variation [10, 23,24]. Fig. 6 shows central microstructure of the cross-sections, where is possible to note same constituents’ types in all coatings. Difference is concentrated in the distribution, morphology and size of these. Micro structure was established through a dendritic formation, whose distri bution uniformity varied according to dilution, where, compared to C1; C2 and C3 present a matrix with smaller dendritic structures. EDX analysis (Fig. 7) from the microstructure characteristics, indicated: point
� Wear regime defined by the Ha/H ratio (where Ha and H are the abrasive and material hardness, respectively) of Khruschov (1957) [18]. The quartz mineral hardness (1160 HV1) was considered as the abrasive hardness of the sand used. � Volumetric loss: measured by weighting the samples made before and after the tribological tests. From the mass loss values, using Eq. (1), these were converted to volumetric loss. (1)
Q ¼ (ml/ρ) * 1000 3
Where: Q ¼ Volumetric loss (mm ); ml ¼ mass loss (g) and ρ ¼ material density (g=cm3 ). � Evaluation of the dimensional wear coefficient per unit sliding dis tance (k) through the Archard model (1953) [19]: from the volu metric loss measured previously, k coefficient was estimated by the ratio presented in Eq. (2). (2)
k ¼ Q / (LN * D) 3
Where: k ¼ Archard’s wear coefficient [mm /(Nm)]; LN ¼ normal load and D ¼ sliding distance. � Wear surface analysis: SEM technique was used to evaluate the wear tracks formed during the tribological tests. The main purpose of this application was to identify the wear micromechanisms action. 3. Results and discussion 3.1. Microstructure The deposition parameters selection methodology resulted in coat ings with different characteristics. The top surfaces showed uniformity, with absence of voids and other discontinuities. Beads width presented a reduction proportional to the scanning speed (V) increase, even increasing laser power (P). Cross-sections geometric characteristics were very similar to the planned performed. Fig. 3 shows the coatings cross sections visualized by SEM. Fig. 3 reveals regular geometries, without inclusions and cracks. Some lack of metallurgical adhesion points on the C1 coating (high lighted in red) were observed, which did not compromise its integrity and tribological performance at the top surface. This discontinuity _ appearance is related to the energy density. As the powder feed rate (m) was kept constant, the laser beam power (P) employed in C1 was insufficient to melt the entire amount of powder added in the fusion set, together with enough substrate portion, generating a poorer metallur gical adhesion. Coatings C2 and C3 were deposited with higher scanning speed than C1, higher laser beam power, but with a lower energy input, resulting in melting all the added powder in these cases, preventing these discontinuities appearance. Evaluating the surface of both coatings, cracks presence is visible, which vary in size. The cracks arise due to the high tensile residual stresses resulting from the cooling process. Greater P and V resulted in coatings with higher cracks concentration, indicating a higher cooling rate. The result was verified by analyzing the C3 surface (deposited with higher P and V), which revealed a larger quantitative number of cracks, as shown in Fig. 4. In Fig. 4, it is observed that the cracks are oriented perpendicularly to the deposition direction, reinforcing the harmful effect of the residual tensile stresses resulting from the cooling process. Guo et al. evaluated Ni-Cr-B-Si alloy coatings deposited under similar conditions. The au thors observed cracks of this nature appearance and emphasize that the coatings deposited on preheated substrate to 300 � C did not present any �k et al. analyzed cracks [20]. Deschuyteneer et al. (2015) and Vost�ra Ni-Cr-B-Si coatings with and without WC addition as reinforcement 4
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 3. Coatings cross sections obtained via SEM (30x).
analyzed a Ni-Cr-B-Si alloy coating deposited by laser cladding with different V levels (5, 10, 15 and 20 mm/s) and observed that rein forcement phases size reduced gradually for higher V [8]. The results of these experiments are similar to those obtained in this research. Similar to this paper, Deschuyteneer et al. (2015) highlight the presence of a microstructure composed of large Cr agglomerates quan tities dispersed in a Ni-rich matrix. As for the other alloy elements, good distribution uniformity was observed throughout the matrix [7,21]. In the literature, studies that carried out X-ray diffraction (XRD) analyzes on Ni-Cr-B-Si alloys coatings report several types carbides observation. Guo et al. point out the lack of clarity regarding the phases type present in the matrix of this alloys class [20]. The different microstructure characteristics identified directly reflect mechanical properties of hardness and coatings tribological behavior. 3.2. Microhardness The average microhardness values of coatings are shown in Fig. 8. It is interesting to note that, even employing relatively different deposition
Fig. 4. Cooling cracks in the coating C3 surface.
Fig. 5. Different microstructure regions visualized via MO (200x).
1 (Pt1, dark region of Fig. 6) - Cr carbides (5CrC), indicated by arrows; point 2 (Pt2, gray region of Fig. 6) - cut dendritic arm and point 3 (Pt3, light region of Fig. 6) - eutectic matrix composed predominantly of Ni and columnar dendrites, highlighted in red. In the literature, by EDX analysis, Deschuyteneer et al. (2015) report the similar phases obser vation [21]. Table 3 shows chemical composition of the points shown in Fig. 7. In Fig. 6, lower concentration and size of the reinforcement phases (CrC) in the higher dilution coatings is observed. This result is consistent �k et al. and Hemmati; Ocelík; De Hosson, who noticed a with Vost�ra proportional number and size phases reduction with P increase, causing a decrease in microhardness, due to a higher dilution [22,23]. Chen et al.
5
parameters, resulting in significant dilution, microhardness average was similar. However, it is important to highlight that the most extern region of the sample is less affected by dilution, which may justify this simi larity. The more intern regions (in the depth direction) are most affected by the increase in Fe content, which reduces the coatings mechanical resistance. In the literature, there are a number of researches reporting different microhardness results from the laser cladding deposition parameters variation, either individually or together. Chen et al. evaluated different scanning speeds and observed that higher velocity of the analyzed ranges exceeded by more than 200 HV0.2 previous coating microhard ness, a behavior attributed to lower energy amount [8]. This result is compatible with the one observed in this paper since parameterization applied generated similar energy input. Fig. 8 also shows the correlation between microhardness and dilution, where it is possible to observe an
CrC: Cr carbides. 5
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 6. Coatings central region microstructure obtained via MEV (2000x).
inversely proportional behavior. There was a gradual increase from C1 to C2, followed by a larger steeper between C2 and C3. This result is intrinsically linked to its relationship with chemical composition and
coatings microstructure. Regardless of the deposition technique, microstructure exerts influ ence on the microhardness. In the literature, an interesting example is
Fig. 7. EDX of the different microstructure phases identified via MEV. 6
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Table 3 Chemical composition of the different microstructure regions identified. Chemical Element
Point 1 (wt.%)
Point 2 (wt.%)
Point 3 (wt.%)
Carbon Silicon Chromium Iron Nickel
45.0 0.4 45.8 0.9 5.9
43.2 1.8 4.9 3.6 43.8
42.6 1.0 4.4 2.1 47.8
Fig. 10. Microhardness/wear coefficient correlation.
phases as it approaches the substrate. García et al. employed the laser cladding process with deposition parameters and a commercial Ni-Cr-B-Si alloy like this work, with and without WC addition. In the base alloy microhardness profile analysis, analogous to this one, the authors observed higher values in the regions close to the coatings top (about 500 HV1), a result that was also attrib uted to a higher CrC phases concentration in this region [26]. Kaiming et al. evaluated the Ni45 alloy resistance, which chemical composition is similar to this work. Deposition parameters used by the authors resulted in a higher microhardness in the HAZ (600 HV0.5), while in the coating (500 HV0.5). This result has been attributed to the occurrence of quenching in this region during coating deposition [27]. Observing at the average surface microhardness, the difference be tween coatings is relatively small. This result is probably related to similar microstructure characteristics in this region. As the same phases’ types were identified and the microindentation hardness test was quite punctual, no significant difference was identified. This result also asserts the high integrity of coatings. Higher variation was observed in the microhardness profiles, where elevated thermal gradient processed coatings presented inferior performance, due to the higher HAZ, greater planar microstructure band (with a higher Fe content and lower hard ness) and lower concentration and size of reinforcement phases dispersed in its matrix.
Fig. 8. Microhardness/dilution relation.
3.3. Wear results
Fig. 9. Microhardness profiles.
3.3.1. Volumetric loss Table 4 shows abrasive wear regime types identified through the Khruschov Ha/H ratio (1957) [18]. In the coatings tested via dry san d/rubber wheel test, the ratio presented a value higher than the limit of 1.2. This result indicates that all coatings were subjected to a severe wear regime. As the difference between the surface microhardness average was small, the ratio presented close values. With the results of volumetric loss that is proportional to the k wear coefficient, it was observed:
presented by Da Silva; D’Oliveira (2015), who compared microstructure and microhardness of Ni-Cr-B-Si coatings deposited by PTA and laser cladding. The authors concluded that in both techniques, the dilution significantly impacted the microstructure and coatings microhardness. The different cooling rate between the two techniques directly modified the microstructure [24]. Fig. 9 shows average microhardness profiles. The coatings presented an average higher to 550 HV1, value almost four times higher to sub strate (150 HV1). A considerable decrease is observed near the sub strate/coating interface, including diluted region and HAZ, respectively. Microhardness profiles followed microstructure variation, where three levels were verified, corresponding to the three microstructure regions. Microhardness increased proportionally to the CrC concentration elevation towards coatings top. In contrast, a proportional decrease of this was observed in relation to reduction in the amount and size of CrC
� In the first 5 min of test, the wear is high, which is probably attrib uted to the following factors: contact area increases between the specimens; vibrations at the moment of wear interface establishment and oxide layer remove. This last phenomenon was an issue explored by Bressan; Hesse; Silva Jr. (2001), who highlight the same as being one of the main oscillations causes in the initial tribological tests phase [28]. Besides to the mentioned characteristics, coatings wear had an additional influence source, which are the cooling cracks present on the sample’s surfaces. During the tests, these acted as anchoring points, generating coating fragments detachments (sub ject to be discussed in more detail in topic 3.3.3).
Table 4 Wear regime by Khruschov’s ratio. Coating
C1
C2
C3
Ha/H ratio Regime type
2.0 Severe
2.1 Severe
2.2 Severe
7
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 11. Features worn out tracks.
� After the stabilizing step, there is a drop in the volumetric loss for 10 min, indicating the wear transition. This occurs due that valleys left by removed fragments are being filled by abrasive material that is not effectively removing coating material. � A new material detachment occurs close to 15 min of the test, evi denced by another high volumetric loss point. � For 20 min of test, a behavior similar to 10 min was observed, probably related to the same reason. � Between 25 and 30 min, average loss and standard deviation differ ence are small compared to the other pauses performed in the test, indicating the uniform volumetric loss regime establishment. The average accumulated volumetric loss over the 30 min of test showed an inversely proportional correlation to microhardness. C1 showed the lowest volumetric loss, followed by C2 and C3, respectively. The matrix characteristics and reinforcement phases (CrC) were the most significant influence sources. The higher concentration and phase size of CrC showed a positive effect on the microhardness and conse quent resistance to abrasive wear. The correlation analysis of hardness and volumetric loss shows a better performance of the coatings with the highest concentration and size of CrC phases. This result is because these coatings are less affected by the sand action when compared to the matrix, which has a lower hardness. Due to the larger dimensions, these phases tend to stay fixed longer in the matrix, giving it a longer protection time. These charac teristics justify the better performance of C1. In the literature, Deschuyteneer et al. evaluated the abrasion resis tance of a Ni-Cr-B-Si alloy with and without WC addition by the dry €kki; Vouristo (2009) [29], sand/rubber wheel test [7]. Nurminen; Na they analyzed wear resistance of various Metal Matrix Composites (MMC) coatings deposited via laser cladding, through the dry san d/rubber wheel test. Alloy studied by the first author and one of the alloys evaluated by the second author is similar to this work. Both au thors highlight obtention of higher wear resistance in coatings that presented higher concentration and size of reinforcement phases. 3.3.2. Archard wear coefficient Fig. 10 shows an inversely proportional correlation between micro hardness and k wear coefficient. The results also expose the stability of this coefficient. Throughout the test, the wear depth increases and other layers of microstructure are affected. As the k coefficient average remained constant, the microstructure variation along depth was not enough to generate significant changes in the coatings. However, the different microstructure characteristics between the three coating con ditions influenced in the k wear coefficient. Moreover, at C2 and C3, wear level was sufficient to reach the hypoeutetic bands and even the planar region, which have lower resistance in relation to the eutectic range. In
Fig. 12. Wear micromechanisms identified via MEV (1500x).
8
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
Fig. 13. Wear micromechanisms interruption (1500x).
this case, it is noticed that C3 indicate of inability to withstand more severe test conditions since their inverse correlation with microhardness presents practically nonexistent. 3.4. Wear surfaces 3.4.1. Wear micromechanisms In general, volumetric loss and the respective k wear coefficients were more pronounced in the coatings with higher dilution, higher cracks concentration, lower amount and size of CrC phases and conse quent lower microhardness. These results reflected directly on the worn surfaces aspects. Abrasive wear mechanism action was observed through grooves (Fig. 11) formed on the sample’s surfaces. An interesting point is that these grooves presented deeper in the regions between beads, a result that can be related to the greater Fe migration to the top of the coatings through the convection currents resulting from solidification process. The fact of this effect is more pronounced in the higher dilution coatings (higher Fe content) reinforces this justification, indicating that these are also lower microhardness regions. During the tests, sand flow tend re mains trapped at these regions, increasing wear. In the literature, Deschuyteneer et al. report similar behavior in the worn surfaces ana lyzes of their coatings [7]. In this research, wear micromechanisms types identified were com mon to all coatings, and the difference was concentrated in their quantitative and severity action. C1 surface presented less severe wear aspects when compared to the other coatings, a result related mainly to its higher microhardness, which prevents the penetration of abrasives on the surface, reducing material removal. Abrasive wear predominant action was observed through microploughings and microcuttings. There are also microindentations characteristic of the wear test and CrC phases fractured by the abrasive particles action. In the specimen’s central re gion, all coatings presented the microploughing as predominant wear micromechanism. The microcutting action was inhibited in regions with CrC phases presence, a more prominent aspect was observed in C1. Fig. 12 shows these worn surfaces. In some points, transitions from microploughing to microcutting are observed, going back to the wear micromechanisms transition model of Zum Gahr [30]. By comparing results of Fig. 12 with k wear coefficient, worn surfaces aspects were consistent, being more severe in C2 and C3, which pre sented performance inferior to C1. However, the intensity degree of micromechanisms identified action was relatively contained. In this way, a more detailed analysis was carried out on the influence of CrC phases and cooling cracks present in the coatings.
Fig. 14. Worn surfaces close to the cooling cracks obtained via MEV: C2 (500x) and C3 (1000x).
3.4.2. CrC carbides and wear surfaces The coatings surfaces presented wear tracks arranged in random directions. No CrC phases detachments of the matrix were observed, and its protective action was identified in several regions, more highlighted on the C1 surface. In Fig. 13, regions where CrC phases act as barriers, preventing the abrasive particles advance on the coating surface, are clearly observed. Besides that, there are points where CrC phases could not avoid, however, they made difficult the continued abrasive sand action flow, resulting in the less severe microploughings formation. In the literature, most of the authors that used dry sand/rubber wheel test highlight an increase in the coatings abrasion resistance through carbides concentration increasing [7,13,21,29,31,32]. Similar to behavior observed in this work, the authors highlight as advantages hardness enhancement, which increases the surface resistance to abra sive particles penetration, the restraint effect that these imposes to the wear micromechanisms advancement, besides the reduction of matrix area in direct contact with wear interface. However, increase in the reinforcing phases concentration is only interesting up to a certain limit. In this sense, Chatterjee; Pal evaluated a ratio of wear rate/carbide concentration and describe that from a given point, the excess carbides became harmful, where the number of frac tured phases acting as abrasive particles increased too much [33].
9
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002
3.4.3. Cooling cracks and wear surfaces Coatings with higher cooling cracks concentration showed higher wear. SEM analyses (Fig. 14) revealed coating material detachments in the cracks and regions close to these ones. This behavior was more evident on the C3 surface when compared to C2. On the other hand, C1 did not present this problem, a result that can be attributed to the lower amount and intensity of cooling cracks identified on its surface. The more highlighted cracks concentration in C2 and C3 is due to the com bination between higher P and V used in these that resulted in a higher thermal gradient. In Fig. 14C2, a crack is observed parallel to the deposition and dry sand/rubber wheel test direction, respectively, due to the tensile stresses resulting from the cooling process. It is possible to observe coating material detachment in the cracks and in nearby regions. Besides that, above the removed region, the eminence of a new coating material detachment is visible, which would increase the wear tracks formed and respective side effects of this phenomenon. This behavior tends to occur cyclically as the test proceeds, generating larger and tending to a cascade effect. Fig. 14C3 presents same behavior, however, in a more severe and deconcentrated way, where along the length of the entire crack, there are detachments, making it even difficult to identify regions that only show cracks. Considering the Fig. 14C2 and Fig. 14C3, it is possible to infer that wear identified on C2 surface is more specific and severe than on C3. However, in C2, cracks collapsed in a restricted region, while in the second, effect covered much of the surface. Higher dilution and conse quent lower microhardness of C3, together with its higher cracks con centration, are consistent with this result. In the literature, works that employ most varied deposition process types, report similar problems. For example, Colaço; Maranho who employed the flux-cored arc welding coating process, as also Natarajan et al. who used laser cladding tech nique report that microploughings formed by the abrasive sand sliding removed large coating material fragments in the regions near cooling cracks, increasing wear [13,32]. Although of being different cases with a similar purpose, the key problem indicated is basically the same, which also extends to the respective work. In theory, material detachment close to the cracks can be charac terized as brittle rupture. However, in this work, it was not possible to confirm microcracking action associated with this behavior, since these cracks origin is known. The large coating fragments loss, caused by material breakage at the cracking limit is also indicated as a key factor in the work of Houdkov� a et al. [12]. Another point observed on both coatings surface is some dark stains, from the abrasive sand accumulation, where a portion of these was fractured and embedded in the surfaces of the coating, acting as a pro tective film, contributing to wear reduction. These stains origin was confirmed by EDX carried out in the region, which indicated a high content of Si (20%) and oxygen O (14%), silica constituent elements (SiO2). Coronado Marin observed similar stains on the AISI 1045 steel surface after performing dry sand/rubber wheel test, attributing them to the same justification [31]. Evaluating the general worn surfaces aspects, the positive influence of the CrC phases presence and the highly cooling cracks damaging ef fects, it is possible to infer that the microploughing action associated to the brittle material detachment parallel to the cooling cracks were predominant causes of the observed wear levels. Besides that, micro cuttings and microindentations presence can also be indicated as responsible for abrasive wear rate.
�
�
�
�
integrity and tribological performance. According to the powder feed rate, parameters used to this coating gave it a lower dilution in comparison to the others. Regions with highest cooling cracks observed in C2 and C3 were attributed to the higher thermal gradient and these perpendicular arrangement defects confirmed the damaging effect of the tensile stresses resulting from cooling process. The microstructure formed by the chemical composition gradient and G/R ratio was constituted by CrC phases dispersed in a Ni den dritic eutectic matrix, whose size and distribution varied according to the dilution. With the surface and cross sections analyzes of C1, was identified a higher microhardness average, that is due to modifying mechanical and metallurgical properties. Microhardness profile showed three different levels, corresponding to the three different microstructure regions. Establishing a magnitude order from higher to lower abrasive wear resistance, coatings can be classified: C1, C2 e C3, a result that can be attributed mainly to microstructure and microhardness characteris tics, which presented inversely proportional with the results of volumetric loss and k wear coefficient. In the worn surfaces analyzes, a considerable wear micromechanisms part were interrupted for the restraint effect formed by CrC phases. It is possible to infer that the regions with higher material volume removed, was due to cooling cracks. In general, it was observed a better tribological performance of C1 in comparison to the other coatings, within the analyzed context, highlighting it to the most suitable for more severe wear conditions applications. However, depending on the solicitation degree, C2 and C3 are also reproducible, since the higher concentration and severity of the cooling cracks identified in these can be reduced through substrate preheating treatments.
Declaration of competing interest None. Acknowledgements The Federal University of Santa Catarina (UFSC, BRA) and SATC College (SATC, BRA) by scientific technical support. The Coordination for the Improvement of Higher Education Personnel (CAPES, BRA) and National Council for Scientific and Technological Development (CNPq, BRA) for financial support. References [1] H€ arkisaari P. Wear and friction effects on energy consumption in the mining industry. Master of science thesis. Tampere: Tampere University of Technology; 2015. [2] Holmberg K, Erdemir A. Influence of tribology on global energy consumption, costs and emissions. Friction 2017;5:263–84. https://doi.org/10.1007/s40544-0170183-5. [3] Holmberg K, Kivikyt€ o-Reponen P, H€ arkisaari P, Valtonen K, Erdemir A. Global energy consumption due to friction and wear in the mining industry. Tribol Int 2017;115:116–39. https://doi.org/10.1016/j.triboint.2017.05.010. [4] Toyserkani E, Corbin S, Khajepour A. Laser cladding. Boca Raton (USA): CRC Press LLC; 2005. [5] Zhong M, Liu W. Laser surface cladding: the state of the art and challenges. J Mech Eng Sci 2010;224:1041–60. https://doi.org/10.1243/09544062JMES1782. [6] Davim JP. Laser in manufacturing. Hoboken: John Wiley & Sons; 2013. [7] Deschuyteneer D, Petit F, Gonon M, Cambier F. Influence of large particle size–up to 1.2 mm–and morphology on wear resistance in NiCrBSi/WC laser cladded composite coatings. Surf Coat Technol 2017;311:365–73. https://doi.org/ 10.1016/j.surfcoat.2016.12.110. [8] Chen J, Li J, Song R, Bai LL, Shao JZ, Qu CC. Effect of the scanning speed on microstructural evolution and wear behaviors of laser cladding NiCrBSi composite coatings. Opt Laser Technol 2015;72:86–99. https://doi.org/10.1016/j. optlastec.2015.03.015. [9] Luo X, LI J, LI GJ. Effect of NiCrBSi content on microstructural evolution, cracking susceptibility and wear behaviors of laser cladding WC/Ni–NiCrBSi composite coatings. J Alloy Comp 2015;626:102–11. https://doi.org/10.1016/j. jallcom.2014.11.161.
4. Conclusions From the results obtained, it was possible to reach some conclusions: � The coatings presented a good surface finish for shape quality. Points with lack of metallurgical adhesion identified in C1, attributed to the insufficient energy density parameter did not compromise its 10
J.M.S. Sousa et al.
Tribology International 143 (2020) 106002 � anek C. The influence of process �k M, Houdkov� [22] Vost�ra a S, Bystrianský M, Ces� parameters on structure and abrasive wear resistance of laser clad WC-NiCrBSi coatings. Mater Res Express 2018;5. https://doi.org/10.1088/2053-1591/aad859. [23] Hemmati I, Ocelík V, De Hosson JTM. Dilution effects in laser cladding of Ni–Cr–B–Si–C hardfacing alloys. Mater Lett 2012;84:69–72. https://doi.org/ 10.1016/j.matlet.2012.06.054. [24] Da Silva LJ, D’Oliveira ASCM. Liga NiCrSiBC: microestrutura e Dureza de Revestimentos Processados a Arco e a Laser. Soldagem & Inspeç~ ao. 2015;20:39–47. https://doi.org/10.1590/0104-9224/SI2001.05. [25] Kou S. Welding metallurgy. New Jersey: John Wiley & Sons; 2003. [26] García A, Fern� andez MR, Cuetos JM, Gonz� alez R, Ortiz A, Cadenas M. Study of the sliding wear and friction behavior of WC þ NiCrBSi laser cladding coatings as a function of actual concentration of WC reinforcement particles in ball-on-disk test. Tribol Lett 2016;63:41. https://doi.org/10.1007/s11249-016-0734-3. [27] Kaiming W, Yulong L, Hanguang F, Yongping L, Zhenqing S, Pengfei M. A study of laser cladding NiCrBSi/Mo composite coatings. Surf Eng 2016;34:267–75. https:// doi.org/10.1080/02670844.2016.1259096. [28] Bressan J, Hesse R, Silva Jr E. Abrasive wear behavior of high speed steel and hard metal coated with TiAlN and TiCN. Wear 2001;250:561–8. https://doi.org/ 10.1016/S0043-1648(01)00638-X. [29] Nurminen J, N€ akki J, Vuoristo P. Microstructure and properties of hard and wear resistant MMC coatings deposited by laser cladding. Int J Refract Metals Hard Mater 2009;27:472–8. https://doi.org/10.1016/j.ijrmhm.2008.10.008. [30] Zum Gahr K-H. Microstructure and wear of materials, tribology series. New York: Elsevier Science; 1987. [31] Coronado Marin JJ. Efeito do tamanho do abrasivo no desgaste de metais. Tese de Doutorado. S~ ao Paulo: Universidade de S~ ao Paulo; 2010. [32] Colaço FHG, Maranho O. Avaliaç~ ao da perda de massa de revestimento duro depositado por soldagem com arame tubular de liga FeCrC-Ti. Soldagem & Inspeç~ ao. 2014;19:58–68. https://doi.org/10.1590/S0104-92242014000100008. [33] Chatterjee S, Pal T. Wear behaviour of hardfacing deposits on cast iron. Wear 2003; 255:417–25. https://doi.org/10.1016/S0043-1648(03)00101-7.
[10] Da Silva LJ, D’OLIVEIRA ASCM. NiCrSiBC coatings: effect of dilution on microstructure and high temperature tribological behavior. Wear 2016;350: 130–40. https://doi.org/10.1016/j.wear.2016.01.015. [11] Miguel J, Guilemany J, Vizcaino S. Tribological study of NiCrBSi coating obtained by different processes. Tribol Int 2003;36:181–7. https://doi.org/10.1016/S0301679X(02)00144-5. � Smazalov� �k M, Schubert J. Properties of NiCrBSi coating, as [12] Houdkov� a S, a E, Vost�ra sprayed and remelted by different technologies. Surf Coat Technol 2014;253: 14–26. https://doi.org/10.1016/j.surfcoat.2014.05.009. [13] Natarajan S, Anand EE, Akhilesh KS, Rajagopal A, Nambiar PP. Effect of graphite addition on the microstructure, hardness and abrasive wear behavior of plasma sprayed NiCrBSi coatings. Mater Chem Phys 2016;175:100–6. https://doi.org/ 10.1016/j.matchemphys.2016.02.076. [14] Yin B, Zhou HD, Chen JM, Yan FY. Microstructure and high temperature sliding wear behaviour of laser remelted NiCrBSi/35(WC-12Ni) composite coating. Surf Eng 2011;27:458–63. https://doi.org/10.1179/026708410X12687356948517. [15] Ocelík V, Nenadl O, Palavra A, De Hosson JThM. On the geometry of coating layers formed by overlap. Surf Coat Technol 2014;242:54–61. https://doi.org/10.1016/j. surfcoat.2014.01.018. [16] ASTM-E92. American society for testing and materials: standard test methods for Vickers hardness and knoop hardness of metallic materials. 2017. [17] ASTM-G65. American society for testing and materials: standard test method for measuring abrasion using the dry sand/rubber wheel apparatus. 2016. [18] Khruschov MM. Resistance of metals to wear by abrasion, as related to hardness. Conf. Lubrication and Wear 1957:655. [19] Archard J. Contact and rubbing of flat surfaces. J Appl Phys 1953;24:981–8. [20] Guo C, Zhou J, Chen J, Zhao J, Yu Y, Zhou H. High temperature wear resistance of laser cladding NiCrBSi and NiCrBSi/WC-Ni composite coatings. Wear 2011;270: 492–8. https://doi.org/10.1016/j.wear.2011.01.003. [21] Deschuyteneer D, Petit F, Gonon M, Cambier F. Processing and Characterization of laser clad NiCrBSi/WC composite coatings - influence of microstructure on hardness and wear. Surf Coat Technol 2015;283:162–71. https://doi.org/10.1016/ j.surfcoat.2015.10.055.
11