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nanocrystalline composite coating by plasma spraying: Effect of heat input on morphology, phase evolution and mechanical properties
Accepted Manuscript Fe-based amorphous/nanocrystalline composite coating by plasma spraying: Effect of heat input on morphology, phase evolution and mechanical properties Anil Kumar, Rahul Kumar, Pavan Bijalwan, Monojit Dutta, Atanu Banerjee, Tapas Laha PII:
S0925-8388(18)33253-5
DOI:
10.1016/j.jallcom.2018.09.024
Reference:
JALCOM 47443
To appear in:
Journal of Alloys and Compounds
Received Date: 21 June 2018 Revised Date:
21 August 2018
Accepted Date: 2 September 2018
Please cite this article as: A. Kumar, R. Kumar, P. Bijalwan, M. Dutta, A. Banerjee, T. Laha, Fe-based amorphous/nanocrystalline composite coating by plasma spraying: Effect of heat input on morphology, phase evolution and mechanical properties, Journal of Alloys and Compounds (2018), doi: 10.1016/ j.jallcom.2018.09.024. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT Fe-based amorphous/nanocrystalline composite coating by plasma spraying: Effect of heat input on morphology, phase evolution and mechanical properties
Tapas Laha a,* a
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Anil Kumar a, Rahul Kumar a, Pavan Bijalwan b, Monojit Dutta b, Atanu Banerjee b, and
Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur, 721302, India
Research and Development Division, Tata Steel, Jamshedpur, 831007 India
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b
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*Corresponding author: Tapas Laha; Email: [email protected]; Phone: +91-3222-
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ACCEPTED MANUSCRIPT Abstract Fe-based amorphous/nanocrystalline composite coatings with a lean composition of Fe−2.5Cr−6.7Si−2.5B−0.7C (wt%) were synthesized by atmospheric plasma spraying (APS) onto a mild steel substrate. The effects of plasma power on the morphology and the phase
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content of the coatings were systematically investigated. Denser coatings with better intersplat bonding were obtained at a higher plasma power, which was attributed to higher degree of powder melting. The retention of amorphous phase and formation of various
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nanocrystalline Fe-borides in the amorphous matrix was decided by the variation in plasma
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power, which in turn affected the mechanical properties of the coatings. Increasing plasma power resulted in higher hardness and elastic modulus of the coatings, which is attributed to the compact microstructure of the coatings containing amorphous matrix with nanocrystalline intermetallics (Fe23B6, Fe2B, and/or Fe3B) distributed. The nanoscratch results indicated that the increased plasma power resulted in uniform scratch profiles. Moreover, dry sliding wear
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test showed that both coefficient of friction and wear rate decreased with increasing plasma power. An analytical model was used to correlate mechanical and tribological properties of
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the coatings, which insinuated that the coating prepared at a plasma power of 35.5 kW exhibited significantly high shear strength than other coatings deposited at lower plasma
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power and approximately 3.6 times greater than that of the mild steel substrate.
Keywords: Fe-based coating; Amorphous/nanocrystalline composite; Atmospheric plasma spraying; Microstructural analysis; Nanoindentation; Wear
ACCEPTED MANUSCRIPT 1. Introduction Bulk metallic glasses (BMGs) possess excellent mechanical properties and high corrosion resistance owing to the absence of long-range atomic periodicity and defects, such as grain boundaries, anti-phase boundaries etc. [1-5]. However, industrial applications of BMGs are
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rather restricted due to the difficulty encountered in bulk quantity production of BMGs
attributed to their limited glass forming ability (GFA) along with intrinsic brittleness at room temperature. In contrast, metallic amorphous coatings, which are easier to synthesize in
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comparison with BMG synthesis, attributed to the lower dimension of coatings in one
direction i.e. thickness, have received considerable attention in the recent years because of the
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combination of the excellent hardness, wear resistance and corrosion resistance owing to the amorphous nature and the potential engineering applications as coatings [6-10]. Compared to monolithic metallic glass coatings, glassy coatings of composite nature with crystalline phases distributed in the amorphous matrix have received considerable scientific and
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industrial interest towards further improving the mechanical properties. In recent days, Febased amorphous/nanocrystalline composite coatings deposited by thermal spray technologies (e.g. flame spray, electric arc spray, plasma spray, and high-velocity oxygen
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fuel spray) are being researched as a viable option for long-term surface protection of boilers,
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gas turbines, hydraulic turbine blades etc. [6, 7, 11-17]. In order to obtain amorphous/nanocrystalline composite coatings, atmospheric plasma spraying (APS) is considered as a simple, versatile and efficient coating process in both research community and industries [18-23]. This process provides a very high cooling rate (~107-108 K/s), which favors the formation of amorphous phase. On the contrary, simultaneous particle oxidation and heat build-up due to deposition of successive molten splats cause crystalline phase precipitation in the amorphous matrix [18-22], which significantly influences the properties of the coatings. Zhang et al. synthesized Fe-based
ACCEPTED MANUSCRIPT metallic glassy coating by plasma spraying and observed that lower spraying power resulted in semi-melting of feedstock powders leading to poor splat adhesion and on the other hand, higher spraying power favors the crystalline phase precipitation in the coatings [18]. Other researchers fabricated Fe-based amorphous coatings using powders of different sizes and
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revealed that the finer feedstock powders could improve the cohesion between splats,
decrease the porosity, and form a composite structure (uniformly distributed crystalline
phases in an amorphous matrix), which ultimately improved mechanical and wear properties
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[13]. Koga. et al. reported improvement in wear property of Fe-based amorphous coating with crystalline precipitates such as hard boride phase formation during spray coating process
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[24].
High material cost associated with the elements (e.g. Ni, Hf, Nb, Y etc.), which increase GFA of Fe-based multi-component systems, limits the wide commercial acceptance of these systems [4, 25, 26]. This led to a research drive on synthesizing Fe-based metallic glasses
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using abundant and relatively inexpensive metalloids viz. B, C, Si, and P [4]. The addition of these elements primarily assist in increasing the glass forming ability of multi-component systems and on the other hand, take part in solid solution strengthening along with driving
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precipitation of various hard phases (e.g. carbides or borides) [26].
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As mechanical properties strongly affect the tribological properties of coatings, therefore, understanding the correlation between mechanical and tribological properties of the composite coating becomes very important. Nanoindentation and nanoscratch techniques are very suitable techniques to evaluate mechanical (nanohardness and reduced Young’s modulus) and tribological (lateral force) properties of amorphous/nanocrystalline solid materials [27]. But, investigations are concentrated separately either on nanoindentation or on nanoscratch of coatings and therefore, a conclusive correlation between mechanical and tribological properties of the Fe-based coatings is lacking.
ACCEPTED MANUSCRIPT In the present work, Fe-based amorphous powder with lean composition was used to deposit amorphous/nanocrystalline composite coating on mild steel substrate through atmospheric plasma spraying. The effect of spraying power i.e. the heat input on the microstructure, distribution of different nanocrystalline secondary phases in the amorphous matrix, and its
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effect on the mechanical and tribological properties of the coatings have been correlated
systematically. An analytical model developed by Chen et al. [27] has been used to correlate the mechanical and tribological properties of the coatings. The aim of the present paper is to
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provide theoretical support for defining the suitable spraying power in depositing Fe-based
properties.
2. Experimental procedures 2.1. Coating deposition
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amorphous/nanocrystalline composite coating with better mechanical (hardness and wear)
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Gas atomized Fe-based amorphous powders (Fe-2.5Cr-6.7Si-2.5B-0.7C, wt%) procured from Epson Atmix Corporation, Japan was used as the feedstock for coating deposition. This chosen alloy composition possesses high potential for industrial applications because of its
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high glass forming ability and relatively low cost of the alloying elements. The coatings with a thickness of about 150 ± 40 µm were deposited by an atmospheric plasma spraying system
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(AMTAG Adv. Mat. Tech., Switzerland, MP 200 Pulse Arc 1000) equipped with a 6-axes robotic arm. The feedstock powders were sprayed onto cold rolled close annealed (CRCA) mild steel substrates of dimension 50 mm x 20 mm x 2 mm. Before spraying process, the substrates were grit-blasted using 24 grit silica beads and ultrasonically cleaned with ethyl alcohol. Argon (99.9 % pure) and Hydrogen (99.9 % pure) of constant flow rates were used as primary gas and secondary gas, respectively; whereas the plasma power was varied in the range of 18.6–35.5 kW by keeping plasma current fixed and changing voltage during the
ACCEPTED MANUSCRIPT plasma spraying process. Thus, the plasma power values were obtained as 18.6 kW, 23.7 kW, 29 kW and 35.5 kW while actually aiming for 20 kW, 25 kW, 30 kW and 35 kW, respectively. Higher plasma power than 35 kW was not tried in this work, as higher plasma power leads to high degree of oxidation which promotes devitrification. Besides, smaller
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feedstock powders (mostly in the range of 10 to 30 µm were selected for this work to get the advantage of better inter-splat bonding resulting from uniform melting of lower size powders possessing higher specific surface area. However, lower particles also experience higher in-
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flight temperature which promotes the self-annealing effect after deposition, which could lead to even more devitrification. Therefore, the plasma power was not raised beyond 35 kW,
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keeping in mind the effect of (i) degree of oxidation, (ii) uniform melting and (iii) selfannealing in the deposited coatings. The details of the plasma spraying parameters are shown in Table 1. For convenience, coatings prepared under different plasma spraying conditions are
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denoted as Coating-I, Coating-II, Coating-III, and Coating-IV, respectively.
2.2. Microstructural characterization
The surface morphology and microstructure of the coatings were investigated with a scanning
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electron microscope (SEM, SUPRA 40, Carl Zeiss AG, Germany). The porosity of the
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coatings was evaluated by quantitative image analysis performed on the SEM images using ImageJ k 1.45 software, where area analysis method was applied by calculating the ratio of porosity area to the total area of a particular field of view (FOV). The primary factor to precisely differentiate the gray scale of porosities form that of the rest of the microstructural features is conducting the process of image thresholding meticulously. At first, the primitive SEM images were converted into 32-bit gray scale images. Then, FFT-based band pass filtering was done before performing thresholding operation to minimize the noise which may get introduced into the image near the porosity boundaries after thresholding. A baseline
ACCEPTED MANUSCRIPT thresholding operation was carried out with a carefully selected threshold parameters based on the pixel value histogram [28]. At least ten SEM images with a magnification of 300X were obtained from the polished cross-sections surface of the various coatings and were used subsequently as FOVs for the porosity estimation. The phase formation in the feedstock
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powders and coatings was characterized by X-Ray diffraction (XRD, PANalytical-Empyrean) equipped with a monochromatic Cr-Kα (λ= 0.229 nm) X-ray source in the range of 30° ≤ 2θ ≤ 110°. The phase analysis obtained from XRD results were validated by carrying out
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transmission electron microscopy (FEG-TEM, JEOL, JEM-2100F) on the coatings. For TEM observation, samples of about 0.4 mm thickness were machined out by wire-cut machine
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from the as-sprayed coating surface, perpendicular to the coating deposition direction. These EDM cut samples were thinned down to 30 µm by mechanical grinding, followed by ionbeam milling using a Gatan (model 691) system at an incident angle of 4-6° with a beam
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energy of 4 KeV until perforation.
2.3. Microindentaion and nanoindentation experiments Vickers micro indenter (UHLVMHT - 001, Walter Uhl, Germany) was used with a 25 gf load
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and a dwell time of 10 s to evaluate the microhardness of the sintered samples.
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Microhardness values were obtained from the average of 15 individual measurements performed on the polished top surface of the coatings. Nanoindentation was carried out using a Triboindenter (TI 950, Hysitron Inc., USA) at room temperature with a standard three-sided pyramid Berkovich tip (TI-0039, Hysitron Inc., total included an angle of 142.3º and tip radius of 100 nm) at a peak load of 5000 µN and loading-unloading rate of 500 µN/s. Depthsensing indentation test was carried out on a standard fused silica (Poisson's ratio of 0.17) specimen beforehand for the calibration of the indenter tip area function. The average nanohardness and reduced modulus reported in this paper represent the average of 100
ACCEPTED MANUSCRIPT indents (10 X 10 matrix). The indenter was held at peak load for 10 s to remove any creep effect. The spacing between indentations was set to be 30 µm to avoid the plastic flow effect of previous indentations on subsequent indentations.
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2.4. Nanoscratch and dry sliding wear tests
Nanoscratch test was conducted under ramp loading with a peak force of 5 N at a loading rate of 125 mN s-1 and indenter velocity of 5 µm s-1 over a scratch length of 200 µm. Dry sliding
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wear test on the coatings was conducted by using a pin-on-disc tribometer (TE97 Friction and Wear Demonstrator, Phoenix Tribology Ltd., England) to evaluate the tribological properties.
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The coatings as pin were held against the counter face of a rotating alumina disc with wear track diameter 25 mm. Before tests, the coatings were sequentially ground with 600, 1000, 1200, 1500, and 2000 mesh SiC papers. The polished coatings were ultrasonically cleaned in acetone for 5 min and then dried in air. The sliding tests were conducted at room temperature
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with a sliding distance of 235 m, a load of 10 N, and a constant speed of 0.26 m s−1. The wear rate was calculated by using equation, W = V/SF, where W is the wear rate in mm3/Nm, V is the wear volume loss in mm3, S is the total sliding distance in m, and F is the applied load in
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N. At least three repeat scratch tests were carried out for each coating sample under
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predefined test conditions.
3. Results and discussions 3.1. Morphology and phase structure of feedstock powders Fig. 1(a) shows the SEM micrograph of the Fe-based gas-atomized powders used as feedstock during plasma spraying in the present study. Most of the powder particles are of spherical morphology and exhibit smooth surface, which conforms to the morphological features of powders synthesized by gas atomization. It can be observed from the bar chart of
ACCEPTED MANUSCRIPT particle size distribution, shown in Fig. 1(b) that the particle size varies in the size range of 5 to 55 µm, however most of the particles are of 10-30 µm. Spherical morphology of powders with a smooth surface and this particular size range are indicative of good flowability during the thermal spraying process. It was also found that some of the larger particles were capped
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with smaller ones, as shown by arrows in Fig. 1(a). When particles with different sizes collide to each other in the gas turbulence during atomization process, then smaller particles having higher solidification rate compare to the larger one, would tend to adhere to the molten
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surface of the larger particles [29]. Fig. 1(c) presents the XRD pattern of the feedstock
powders. Presence of a broad halo peak appearing in 2θ range of 50° - 90° without any
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crystalline peak indicates that the powders were amorphous in nature. The formation of amorphous phase is attributed to the high GFA of the Fe−2.5Cr−6.7Si−2.5B−0.7C (wt.%) alloy and fairly high solidification rate during gas atomization of the powders in an inert
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atmosphere.
3.2. Microstructure of as-sprayed coatings
Top surface SEM images of the as-sprayed coatings deposited with varying plasma power
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(18.6, 23.7, 29 and 35.5 kW) are presented in Fig. 2(a-d) to show the splat morphology after
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deposition of the coatings; whereas the polished surface images in Fig. 2(e-h) depict the internal microstructure. It can be clearly seen from Fig. 2(a) that Coating-I, which was deposited at a low spraying power (18.6 kW, Table 1), contained higher extent of partially molten particles on the unpolished coating surface. The corresponding image (Fig. 2e) of the polished surface exhibits high amount of large pores. As the spraying power was gradually increased from 18.6 kW to 35.5 kW, the amount of partially molten particles reduced (Fig. 2b-d). Higher degree of powder melting in case of maximum plasma power of 35.5 kW is evident from the corresponding Fig. 2d. It can also be observed from the micrographs of the
ACCEPTED MANUSCRIPT polished coating surface (Fig. 2e-h) that the amount and size of porosity were gradually reduced as the plasma power increased. The cross-sectional morphology of the various coatings is shown in Fig. 3. All of the four coatings with thickness of around 150 ± 40 µm were well-bonded to the mild steel substrate and the inter-splat bonding increased gradually
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from Coating-I to Coating-IV (Fig. 3a-d), as the plasma power increased. The porosity of coatings tends to decline from Coating-I to Coating-IV and the porosity vol.% were 14.8%, 12.6%, 9.1%, and 4.8% for Coating-I, Coating-II, Coating-III, and Coating-IV, respectively
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(as reported in Table 2). The low spraying power was inadequate to completely melt the
feedstock powders and therefore led to deposition of poorly adhered splats which resulted in
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higher porosity in Coating-I (Fig. 3a). The Coating-ІV possessed the lowest because of welladhered splats (Fig. 3d). The dense structure of this coating is mainly ascribed to the higher degree of melting leading to better spreading of splats upon impact on the substrate and previously deposited splat layers [13, 21, 30]. Higher spraying power provides more enthalpy
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to accelerate the particle melt and reduces the surface tension of droplets; and therefore, the droplets flatten well, leading to the formation of a dense coating structure [19]. Besides, at higher spraying power, gas expands fast and leads to high speed plasma flow and hence
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drives the molten particles to impact intensely on substrate and to form denser coating [21].
3.3. Phase composition of the coatings Fig. 4 shows XRD patterns of the different coatings obtained at different spraying powers. XRD patterns of all the four coatings exhibit a broad halo peak appearing in 2θ range of 55– 85° with some low-intensity crystalline peaks. Thus, the XRD results insinuate synthesis of amorphous/nanocrystalline composite coatings. Comparing the XRD patterns of the feedstock powders (Fig. 1b) and the coatings (Fig. 4), it is evident that precipitation of crystalline phases occurred during the coating deposition process. The crystalline peaks
ACCEPTED MANUSCRIPT appeared in the XRD patterns of the coatings are associated with α-Fe (PDF No. #00-0060696), Fe23B6 (PDF No. #00-047-1332) and B13C2 (PDF No. #00-026-0233) phases. The reasons behind crystalline phase formation can be attributed to the following factors. Firstly, depending on the heat input and associated degree of melting caused by increase in plasma
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power, certain amount of crystallization took place during splat solidification. Secondly, a certain amount of heat build-up took place due to deposition of successive molten splats and release of latent heat of fusion during solidification of these splats [19, 21, 30]. Besides, there
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was a probability of powder oxidation to some extent during the coating deposition attributed to high-temperature processing in atmospheric environment which may lead to the faster
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crystallization during solidification [21, 26]. An et al. also reported increased crystallization in Fe-based metallic glass system with higher degree of oxidation [21]. The amorphous phase content of the samples was estimated from the area ratio of crystalline peaks to amorphous hump [31, 32]. Based on this calculation, retained amorphous phase fraction in Coating-I,
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Coating-II, Coating-III, and Coating-IV were 86 %, 83.6%, 82.8%, and 81%, respectively. This indicates that the fraction of amorphous phase in coatings decreased with increase in plasma arc power. Although there is a trend of increasing crystallinity with increasing plasma
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power, however, the variation is not very significant. The higher amount of heat input in case
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of higher plasma power deposition would lead to more degree of melting (evident from Fig. 2), higher amount of heat build-up in the deposited splats and would have higher probability of oxide formation. All these factors contributed in higher amount of crystallinity in coatings deposited with higher plasma power. It must be mentioned here that phase amount estimation only from XRD patterns is not enough to ensure if a material is completely amorphous. Usually, the identification of small amounts of nanocrystalline phases embedded in high amorphous content matrix could not be clearly identified from the XRD measurements due to the weak intensity peaks of these phases. Moreover, these weak intensity peaks can be
ACCEPTED MANUSCRIPT confused with the background and/or be overlapped by the amorphous phase broad halo peak. The authors intend to carry out exhaustive differential scanning calorimetry on the coatings to find out the amount of retained amorphous phase, where a ribbon of similar composition synthesized by splat melting at a very high wheel speed would be used as the standard fully
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amorphous material.
TEM was employed to understand the phase evolution within the coatings at different plasma power and to validate the composite structure (presence of nanocrystalline and amorphous
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phases) of the coatings. TEM micrographs and corresponding selected area diffraction (SAD) patterns of Coating-I, Coating-II, and Coating-IV are shown in Fig. 5. The TEM micrographs
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(Fig. 5a-c) of these coatings exhibit presence of nanocrystalline phases uniformly distributed in the amorphous matrix and this observation matches well with the corresponding SAD patterns (Fig. 5d-f), which show spotted rings (obtained from nanocrystalline phases) with a diffused background (from amorphous matrix). The TEM images (Fig. 5a-c) reveal that grain
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number density and grain size in the coating increased when the plasma power increased from 18.6 kW to 35.5 kW. This finding confirms that volume fraction of various nanocrystalline phases increased with increasing plasma power and this is consistent with the
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corresponding XRD patterns (Fig.4) showing gradual decrease in hump intensity with
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increasing plasma power.
The HRTEM image obtained from Coating-IV shown in Fig. 6a clearly exhibits the presence of nanocrystalline precipitates embedded in the amorphous matrix which further confirms the composite nature of the produced coating. The selected areas (circles 1 to 4) have been digitally analyzed to obtain the fast Fourier transform (FFT) patterns and reported in Fig. 6(be). Formation of various crystalline phases was envisaged from these FFT patterns. The diffraction spots in Fig. 6(b and c) fit to the α-Fe (1 1 0) plane (d= 2.02 Å) and (2 0 0) plane (d= 1.43 Å) forming an angle of 45° along the [0 0 -1] zone axis, and cubic Fe23B6 (2 4 2)
ACCEPTED MANUSCRIPT plane (d= 2.19 Å) and (4 4 0) plane (d= 1.90 Å) forming an angle of 30° along the [1 -1 1] zone axis, respectively. These d-spacings are well in accord with the d-spacings of various phases which were identified from the XRD patterns presented earlier in Fig. 4. Formation of additional phases, as shown in Fig. 6(d and e), which are tetragonal Fe2B (0 2 0) plane (d=
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2.55 Å) and (1 2 1) plane (d= 2.01 Å) forming an angle of 35° along the [1 0 -1] zone axis and orthorhombic Fe3B showing (1 0 2) plane (d= 2.03 Å) and (2 2 0) plane (d= 2.10 Å) forming an angle of 72° along the [-2 2 1] zone axis, respectively. Formation of similar
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intermetallic phases in Fe-Si-B amorphous systems during devitrification has been reported
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earlier by other researchers [33-36].
3.4. Microindentaion and nanoindentation results of coatings
The mechanical properties including Vickers microhardness, nanohardness, and reduced elastic modulus of the coatings and the mild steel substrate obtained from microindentaion
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and nanoindentation tests are listed in Table 2. All of the coatings exhibited higher microhardness as well as nanohardness than the substrate. The Vickers microhardness of substrate, Coating-I, Coating-II, Coating-III, and Coating-IV were 172 ± 10 HV0.025, 799 ±
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124 HV0.025, 895 ± 123 HV0.025, 954 ± 110 HV0.025, and 993 ± 86 HV0.025, respectively.
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Moreover, the nanohardness of substrate, Coating-I, Coating-II, Coating-III, and Coating-IV were 3.61 ± 0.25 GPa, 9.60 ± 2.30 GPa, 10.51 ± 2.42 GPa, 11.00 ± 2.52 GPa, and 11.80 ± 2.71 GPa, respectively, and the reduced elastic modulus were 184.9 ± 5.6 GPa, 125.6 ± 34.3 GPa, 138.5 ± 38.8 GPa, 154.9 ± 30.8 GPa, and 167.9 ± 31.5 GPa, respectively. It is evident from this data that there is an increase in hardness (both micro and nano) and reduced elastic modulus of the coatings with increase in spraying power. The higher hardness and modulus observed in coatings deposited with higher plasma power can be attributed to (i) the better inter-splat bonding resulting lower porosity of coatings, as shown in Fig. 3 and Table 2, and
ACCEPTED MANUSCRIPT (ii) homogeneously precipitated hard intermetallic phases (Fe2B, Fe3B, and Fe23B6) in the amorphous matrix, (as apparent from the XRD - Fig. 4 and HRTEM - Fig. 6 analysis) which resulted in dispersion strengthening. However, average elastic modulus of the Fe-based metallic glass coatings (125.6 to 167.9 GPa) is considerably lower than mild steel substrate
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(184.9 GPa). This may be due to crystalline (α-Fe) nature of the mild steel substrate and
highly amorphous structure of the coatings. Amorphous alloys have large free volume which increases the inter-atomic distance and weakens the atomic bonding. After crystallization,
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free volume annihilates and results in a relaxed structure with shorter inter-atomic distance, leading to higher hardness and Young's modulus [33, 36-39]. Sun et al. [39] observed that
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partially or fully crystallized structure was showing higher hardness and Young's modulus than the material with a completely amorphous structure. Similar behavior has been observed in present work in terms of hardness and modulus. Average nanohardness increased from 9.6 GPa to 11.8 GPa, and reduced elastic modulus increased 125.6 GPa to 167.9 GPa as
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crystalline content increased at elevated plasma power. Lashgari et al [36] reported that Febased fully amorphous ribbon had reduced elastic modulus of around 160 GPa and increased up to 240 GPa after annealing, due to precipitation of Fe2B and Fe23B6 compounds which are
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extremely brittle. Average reduced elastic modulus of the coatings in the present work
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corroborates well with the earlier reported values [36, 39]. For the coatings having multiphase structure, each individual phase affects the overall properties of the coatings. In this context, grid indentation is a very useful technique for characterization of materials of small volume and to provide quantitative prediction of phase transformations. This two-dimensional surface mapping technique with large arrays of nanoindentation is useful to investigate the properties of individual constituent phases by carrying out statistical analysis of the resulting data (nanohardness and elastic modulus). Resulting nanohardness can be plotted in the form of a histogram of indent number density
ACCEPTED MANUSCRIPT with different hardness range which is associated with different phases present in that material. Based on the nanohardness value reported in previous literatures [13, 36, 39-41], the range of nanohardness data was categorized into four sets viz. 5.0-7.5 GPa, 7.5-9.5 GPa, 9.5-12.5 GPa, and 12.5-20.0 GPa, belonging to (i) mixed nano-sized crystalline α-Fe and Fe-
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based amorphous matrix, (ii) amorphous matrix, (iii) mixed amorphous and nano-sized Febased intermetallic phases, and (iv) nano-sized intermetallics rich regions, respectively. It can be noted here that, the Fe-based amorphous phase has been considered in the first three sets,
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as a wide range of hardness for Fe-based amorphous phase based on presence of various
alloying elements has been reported in various literatures. Recently, Lashgari et al. reported
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the presence of similar nanohardness range for amorphous phase and mixed amorphous and nano-sized borides in Fe-Si-B based alloy system synthesized by melt spinning [36]. Graphical representation of indent number density (%), as shown in Fig. 7, shows the relative fraction of different phases present in the various coatings. From Fig. 7, it can be observed
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that the indent number density (%) with hardness range of 7.5-9.5 GPa, which represents the hardness of amorphous matrix kept on decreasing with increase in plasma power. On the other hand, this percentage for the nano-sized intermetallics rich region with hardness range
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of 12.5-20.0 GPa showed an increasing trend with increasing heat input. It can be commented
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that at elevated spraying power, indenter falling on mixed phase (amorphous and nano-sized intermetallic phases) and nano-sized intermetallics phase has higher probability due to more fraction of nano-sized intermetallics formation. In case of Coating-I, 17% and 30% of nanohardness values were in the range of 5.0-7.5 (mixed nano-sized α-Fe and amorphous matrix) and 7.5-9.5 GPa (amorphous matrix), respectively, in comparison to 5.4% and 9.8% in Coating-IV. The percentage of nanohardness data (9.5-12.5 GPa) related to the mixed amorphous and intermetallic phase in Coating-I and Coating-IV (45.4% and 48.9%, respectively) remained almost similar. Moreover, coating-ІV exhibited more amount of
ACCEPTED MANUSCRIPT intermetallic phases in comparison to Coating-I as 35.9% of hardness data was in 12.5-20.0 hardness range, in comparison to 7.8% for Coating-I. The variation in nanohardness values can also be justified from the TEM analysis. Further, presence of the different phases can be validated by the typical load-displacement curves obtained from grid nanoindentation test as
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plotted in Fig. 8(a). The depth of penetration and corresponding nanohardness values for different type of phases has been depicted in Fig. 8(b). Moreover, maximum indentation
depth was approximately 188 nm for mixed α-Fe and amorphous matrix phase region, and
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approximately 106 nm for the intermetallics region at a constant peak force of 5000 µN, as shown in Fig. 8(b). The decrease in the penetration depth could be justified by the fact that
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intermetallic regions show higher nanohardness and stiffness (increase in the slope of unloading portion of the load-displacement curve). Fig. 8(c) shows 3D depth profile images obtained by in-situ scanning probe microscopy (SPM) carried out on four different regions with indentations on the coating surface. The four different penetration depths (127 nm, 110
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nm, 88 nm and 67 nm) represent the four different hardness range regions (mixed nano-sized α-Fe and amorphous matrix, amorphous phase, mixed amorphous and intermetallic phases, and intermetallic phases) as mentioned earlier. Therefore, the variation in indentation depth
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(from 67 to 127 nm, approximately) validates the presence of several phases in the coating.
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Besides hardness and Young’s modulus, nanoindentation can also be used to predict the wear resistance, which is related to the ratio of H/Er of the coatings [33, 42-44]. The values of H/Er ratio obtained from the nanoindentation with a load of 5000 µN are listed in Table 2. It has been found that coatings have higher H/Er (H/Er ~ 0.07) value than the steel substrate (H/Er ~ 0.019). Many authors have suggested that higher value of the ratio H/Er, generally indicative of the good wear resistance [13, 33, 42-44]. Mukhopadhyay et al. reported that H/Er ratio more than 0.05 indicates high degree of elastic recovery in the material which is another important parameter for higher wear resistance [44]. From the obtained H/Er values of the
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trend of H/Er values became inconsistent.
3.5. Nanoscratch behavior of the coatings
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the rate of change of H and Er. As the rate of increase in H and Er values were not similar, the
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Fig. 9 shows SEM images of the worn-out surfaces of the various coating after the nano-
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scratch test. The worn surfaces of the coatings indicated that grooves became wider with increasing load in scratch direction and coatings deposited at various spraying power have the different capability to resist the splats detachment. Coating-I and Coating-II which were deposited at lower plasma arc powers, displayed severe splat detachments (Fig. 9a and b), due to the lower degree of melting of the feedstock powders and associated poor inter-splat
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bonding. However, due to higher degree of melting of feedstock powders at higher spraying powers (28.6 and 35.5 kW), Coating-III and Coating-IV (Fig. 9c and d) experienced
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improved inter-splats bonding and thus better scratch resistance with no or a very little splats detachment in comparison with Coating-I and Coating-II. Thus, it could be concluded that
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high spraying power resulted in more stable scratch which resulted in comparatively smoother scratch profile (Fig. 9c and d). Additionally, fracture of coating surface is brittle in nature and therefore it can be concluded that coatings are liable to brittle fracture during the friction process. Presence of various nano-sized intermetallics and their distribution in the amorphous matrix would play important role in dictating the wear behavior of the coatings and this issue should be addressed in future work in this area. To understand the tribological behavior of the coatings, dry sliding wear test was carried out, the results of which are reported later in Section 3.7.
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3.6. Correlation between nanoindentation and nanoscratch results Mechanical properties (viz. nanohardness and reduced elastic modulus) and tribological properties (lateral force and scratch path morphology) of the plasma sprayed coatings were
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greatly affected by the change in the spraying power, as shown by the results reported in
Table 2 and Fig. 9. Therefore, correlation between mechanical and tribological properties of the coatings was required to be investigated. Recently, Chen et al. developed an analytical
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model to correlate nano-scale mechanical properties with tribological properties of coatings,
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1/ 2 Fmax H τ F f = 24.5 hm − 2 2.263 × Er
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which is given in equation (1) [27]. (1)
Where, Ff is lateral force (tribological property) obtained during the scratch test, H is hardness, Er is reduced elastic modulus, and τ is shear strength. The effect of surface
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roughness and material sinking-in/pile up on the real contact area during loading has not been included in the above-proposed model. Based on the lateral force values obtained from the
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nanoscratch tests, and elastic modulus and hardness obtained from nanoindentation tests, the shear strength (τ) of the mild steel substrate and the different coatings was estimated using
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Equation 1. The shear strength was found to be 188 ± 12 MPa for the mild steel substrate, 398 ± 56 MPa for Coating-I, 433 ± 46 MPa for Coating-II, 516 ± 54 MPa for Coating-III, and 670 ± 60 MPa for Coating-IV, as listed in Table 2. Coating-IV has significantly high shear strength than other coatings and approximately 3.6 times greater than the mild steel substrate. It can be concluded that higher scratch resistance of the coatings resulted due to the higher shear strength of the respective coatings. Higher shear strength of the coatings leads to higher lateral force during scratch of the coatings, which finally results in higher scratch resistance.
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3.7. Tribological behavior of coatings accessed by dry sliding wear test Fig. 10 depicts the results obtained from pin-on-disc based dry sliding wear tests carried out on the coatings and mild steel substrate at a load of 10 N and sliding velocity of 0.26 m/s.
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Fig. 10(a) illustrates the variation in coefficient of friction (COF) of coatings and substrate as a function of sliding distance. The COF values of the mild steel substrate, Coating-I, CoatingII, Coating-III, and Coating-IV were 0.47 ± 0.06, 0.26 ± 0.05, 0.20 ± 0.02, 0.18 ± 0.03, and
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0.15 ± 0.02, respectively, as reported in Table 2. All coatings showed lower COF than that of the substrate, owing to the higher hardness of the coatings as reported in Section 3.4.
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Moreover, the average COF of the coatings slightly decreased with increasing plasma power as mentioned above. Relatively higher fluctuations in COF curves and higher value of COF were observed for Coating-I compared to other coatings. This was mainly due to its poorly bonded splats (Fig. 3a) with relatively bigger pores (Fig. 2e). The lowest coefficient of
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friction for Coating-IV can be attributed to the high hardness which resulted from the tightly stacked splats (Fig. 3d) and precipitation of higher fraction of very hard Fe-borides phases, as was shown earlier in Fig. 8.
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Fig. 10(b) shows the wear rate of the mild steel substrate and the various coatings. The wear rate of the substrate, Coating-I, Coating-II, Coating-III, and Coating-IV were 37.4 ± 0.9, 16.1
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± 2.2, 10.2 ± 1.5, 8.2 ± 0.8, and 3.9 ± 1.6 mm3/Nm, respectively. The substrate showed the maximum wear rate compared to that of the coatings. A decreasing trend of wear rate of the coatings was observed with increasing plasma arc power. Moreover, the wear rate of the Coating-IV (3.9 mm3/Nm) decreased approximately one-tenth of that of the substrate (37.4 ± 0.9 mm3/Nm), which confirms that the wear resistance of Coating-IV was much superior to that of the substrate. Higher value of H/Er ratio, as discussed in Section 3.5 also supported the fact of higher wear resistance of the coatings. Thus, the deposition of Fe-based
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4. Conclusions
Fe-based amorphous/nanocrystalline composite coatings were developed on a mild steel substrate by atmospheric plasma spraying (APS) of gas-atomized amorphous powders
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(Fe−2.5Cr−6.7Si−2.5B−0.7C, wt%) with varying plasma power. The coatings primarily
consisted of amorphous phase, which was due to the high glass forming ability (GFA) of the
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selected multi-component system and the high cooling rate of APS technique. Higher plasma power resulted in higher degree of melting causing better inter-splat bonding and precipitation of higher fraction of nanocrystalline intermetallics (Fe23B6, Fe2B, and Fe3B) in the amorphous matrix. The nanoscratch results indicated that the increased plasma power
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resulted in uniform scratch profiles due to well-bonded splats. Moreover, dry sliding wear test showed that both coefficient of friction and wear rate of the coatings decreased with increasing plasma power. In the present study, the highest nanohardness of 11.8 ± 2.71 GPa
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and the minimum wear rate of 3.9 ± 1.6 mm3/Nm was achieved for the coating deposited at
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the plasma power of 35.5 kW. An analytical model which was used to correlate the mechanical and tribological properties of the coatings predicted that the coating with higher plasma power has the higher shear strength which ultimately resulted in higher scratch resistance. The present results show that Fe-based amorphous/nanocrystalline composite coatings have great potential for the industrial applications in the areas that involve high wear.
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Acknowledgements The author, T. Laha thankfully acknowledges the financial support obtained from Research
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and Development Division of Tata Steel, India.
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ACCEPTED MANUSCRIPT Table 1. Spray parameters for Fe-based amorphous/nanocrystalline composite coatings deposited on mild steel substrates by atmospheric plasma spraying Coatings Plasma spray parameters Coating-І
Coating-ІІ
Coating-ІІІ 60
Secondary gas (H2) flow rate (L/min)
10
Powder feed rate (g/min)
30
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Primary gas (Ar) flow rate (L/min)
130
Carrier gas (Ar) flow rate (L/min) Current (A) Voltage (V)
3
200
300
400
500
93
79
72.5
71
18.6
23.7
29
35.5
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Plasma power (kW)
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Coating thickness (µm)
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Mild steel substrate
Coating-I
Coating-II
Coating-III
Coating-IV
Porosity (vol.%)
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14.8 ± 2.6
12.6 ± 1.7
9.1 ± 0.8
4.8 ± 0.6
Vickers microhardness
172 ± 10
799 ± 124
895 ± 123
954 ± 110
993 ± 86
Nanohardness (GPa)
3.61 ± 0.25
9.60 ± 2.30
10.00 ± 2.42
11.00 ± 2.52
Reduced elastic modulus (GPa)
184.9 ± 5.6
125.6 ± 34.3
148.3 ± 38.8
149.5 ± 30.8
167.9 ± 31.5
Wear property parameter, H/Er
0.019 ± 0.001
0.076 ± 0.025
0.067 ± 0.033
0.074 ± 0.049
0.071 ± 0.010
Shear strength (MPa)
188 ± 12
398 ± 56
433 ± 46
516 ± 54
670 ± 60
Coefficient of friction (Dry Sliding wear)
0.47 ± 0.06
0.26 ± 0.05
0.20 ± 0.02
0.18 ± 0.03
0.15 ± 0.02
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(HV0.025)
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Fig. 2. SEM micrographs of as-sprayed (a-d) and polished (e-h) top surface morphology of the plasma sprayed coatings deposited with varying plasma power (18.6, 23.7, 29 and 35.5 kW, showing the variation in degree of melting and associated porosity in the various
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Fig. 3. SEM images of cross-sections of as-sprayed iron-based amorphous/nanocrystalline
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coatings deposited with plasma power of (a) 18.6 kW, (b) 23.7 kW, (c) 29 kW and (d) 35.5 kW, showing the variation in morphology.
Fig. 4. XRD patterns of the various plasma sprayed coatings deposited with varying plasma power showing variation in crystallization.
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Fig. 5. TEM images of Coating-I, Coating-II, Coating-ІIІ, and Coating-ІV (a - c) and the respective SAED patterns (d - f).
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Fig. 6. High-resolution transmission electron microscopy (HRTEM) results of Coating-IV: (a) HRTEM image confirming the uniform distribution of nano-sized grains embedded in
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amorphous matrix, (b-e) FFT patterns obtained from the different locations (viz. 1, 2, 3 and 4) in Fig. a corresponding to various crystalline phases i.e. α-Fe, Fe23B6, Fe3B, and Fe2B phases, respectively.
Fig. 7. Bar graph showing the indent number density (%) of the plasma sprayed coatings in four different nanohardness range, viz. 5.0-7.5 GPa, 7.5-9.5 GPa, 9.5-12.5 GPa and 12.5-20.0 GPa, which belong to mixed nano-sized α-Fe and amorphous matrix, amorphous matrix,
ACCEPTED MANUSCRIPT mixed amorphous and nano-sized intermetallics and nano-sized intermetallics rich regions, respectively. Fig. 8. (a) Load vs. penetration depth plots obtained from nanoindentation test corresponding to different phases (amorphous and nano-sized crystalline) present in the Fe-based metallic
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glass composite coatings, (b) the depth of penetration and corresponding nanohardness
obtained from the different phases of Coating-IV deposited at 35.5 kW plasma power and (c) 3D depth profile images obtained by in-situ scanning probe microscopy (SPM) carried out on
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four different indentation regions.
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Fig. 9. SEM images of the scratched surface of the different coatings after ramping load nano-scratch tests at 5 N peak load and 5 µm/s scratch velocity: (a) Coating-І, (b) Coating-ІI, (c) Coating-ІIІ, and (d) Coating-ІV, respectively.
Fig. 10. (a) Coefficient of friction (COF) plotted as a function of sliding distance and (b)
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wear rate obtained from pin-on-disc based dry sliding wear tests carried out on the coatings
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and mild steel substrate at a load of 10 N and sliding velocity of 0.26 m/s.
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ACCEPTED MANUSCRIPT Highlights Fe-based amorphous/nanocrystalline coatings were deposited by plasma spraying.
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Coatings consist of nano-sized α-Fe and Fe-borides distributed in amorphous matrix.
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Higher plasma power resulted in more devitrification and better intersplat bonding.
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High hardness resulted due to denser coating and presence of hard Fe-borides.
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Wear resistance of the coatings increased with increasing plasma power.
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S0925-8388(18)33253-5
DOI:
10.1016/j.jallcom.2018.09.024
Reference:
JALCOM 47443
To appear in:
Journal of Alloys and Compounds
Received Date: 21 June 2018 Revised Date:
21 August 2018
Accepted Date: 2 September 2018
Please cite this article as: A. Kumar, R. Kumar, P. Bijalwan, M. Dutta, A. Banerjee, T. Laha, Fe-based amorphous/nanocrystalline composite coating by plasma spraying: Effect of heat input on morphology, phase evolution and mechanical properties, Journal of Alloys and Compounds (2018), doi: 10.1016/ j.jallcom.2018.09.024. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT Fe-based amorphous/nanocrystalline composite coating by plasma spraying: Effect of heat input on morphology, phase evolution and mechanical properties
Tapas Laha a,* a
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Anil Kumar a, Rahul Kumar a, Pavan Bijalwan b, Monojit Dutta b, Atanu Banerjee b, and
Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur, 721302, India
Research and Development Division, Tata Steel, Jamshedpur, 831007 India
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*Corresponding author: Tapas Laha; Email: [email protected]; Phone: +91-3222-
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ACCEPTED MANUSCRIPT Abstract Fe-based amorphous/nanocrystalline composite coatings with a lean composition of Fe−2.5Cr−6.7Si−2.5B−0.7C (wt%) were synthesized by atmospheric plasma spraying (APS) onto a mild steel substrate. The effects of plasma power on the morphology and the phase
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content of the coatings were systematically investigated. Denser coatings with better intersplat bonding were obtained at a higher plasma power, which was attributed to higher degree of powder melting. The retention of amorphous phase and formation of various
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nanocrystalline Fe-borides in the amorphous matrix was decided by the variation in plasma
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power, which in turn affected the mechanical properties of the coatings. Increasing plasma power resulted in higher hardness and elastic modulus of the coatings, which is attributed to the compact microstructure of the coatings containing amorphous matrix with nanocrystalline intermetallics (Fe23B6, Fe2B, and/or Fe3B) distributed. The nanoscratch results indicated that the increased plasma power resulted in uniform scratch profiles. Moreover, dry sliding wear
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test showed that both coefficient of friction and wear rate decreased with increasing plasma power. An analytical model was used to correlate mechanical and tribological properties of
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the coatings, which insinuated that the coating prepared at a plasma power of 35.5 kW exhibited significantly high shear strength than other coatings deposited at lower plasma
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power and approximately 3.6 times greater than that of the mild steel substrate.
Keywords: Fe-based coating; Amorphous/nanocrystalline composite; Atmospheric plasma spraying; Microstructural analysis; Nanoindentation; Wear
ACCEPTED MANUSCRIPT 1. Introduction Bulk metallic glasses (BMGs) possess excellent mechanical properties and high corrosion resistance owing to the absence of long-range atomic periodicity and defects, such as grain boundaries, anti-phase boundaries etc. [1-5]. However, industrial applications of BMGs are
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rather restricted due to the difficulty encountered in bulk quantity production of BMGs
attributed to their limited glass forming ability (GFA) along with intrinsic brittleness at room temperature. In contrast, metallic amorphous coatings, which are easier to synthesize in
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comparison with BMG synthesis, attributed to the lower dimension of coatings in one
direction i.e. thickness, have received considerable attention in the recent years because of the
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combination of the excellent hardness, wear resistance and corrosion resistance owing to the amorphous nature and the potential engineering applications as coatings [6-10]. Compared to monolithic metallic glass coatings, glassy coatings of composite nature with crystalline phases distributed in the amorphous matrix have received considerable scientific and
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industrial interest towards further improving the mechanical properties. In recent days, Febased amorphous/nanocrystalline composite coatings deposited by thermal spray technologies (e.g. flame spray, electric arc spray, plasma spray, and high-velocity oxygen
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fuel spray) are being researched as a viable option for long-term surface protection of boilers,
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gas turbines, hydraulic turbine blades etc. [6, 7, 11-17]. In order to obtain amorphous/nanocrystalline composite coatings, atmospheric plasma spraying (APS) is considered as a simple, versatile and efficient coating process in both research community and industries [18-23]. This process provides a very high cooling rate (~107-108 K/s), which favors the formation of amorphous phase. On the contrary, simultaneous particle oxidation and heat build-up due to deposition of successive molten splats cause crystalline phase precipitation in the amorphous matrix [18-22], which significantly influences the properties of the coatings. Zhang et al. synthesized Fe-based
ACCEPTED MANUSCRIPT metallic glassy coating by plasma spraying and observed that lower spraying power resulted in semi-melting of feedstock powders leading to poor splat adhesion and on the other hand, higher spraying power favors the crystalline phase precipitation in the coatings [18]. Other researchers fabricated Fe-based amorphous coatings using powders of different sizes and
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revealed that the finer feedstock powders could improve the cohesion between splats,
decrease the porosity, and form a composite structure (uniformly distributed crystalline
phases in an amorphous matrix), which ultimately improved mechanical and wear properties
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[13]. Koga. et al. reported improvement in wear property of Fe-based amorphous coating with crystalline precipitates such as hard boride phase formation during spray coating process
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[24].
High material cost associated with the elements (e.g. Ni, Hf, Nb, Y etc.), which increase GFA of Fe-based multi-component systems, limits the wide commercial acceptance of these systems [4, 25, 26]. This led to a research drive on synthesizing Fe-based metallic glasses
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using abundant and relatively inexpensive metalloids viz. B, C, Si, and P [4]. The addition of these elements primarily assist in increasing the glass forming ability of multi-component systems and on the other hand, take part in solid solution strengthening along with driving
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precipitation of various hard phases (e.g. carbides or borides) [26].
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As mechanical properties strongly affect the tribological properties of coatings, therefore, understanding the correlation between mechanical and tribological properties of the composite coating becomes very important. Nanoindentation and nanoscratch techniques are very suitable techniques to evaluate mechanical (nanohardness and reduced Young’s modulus) and tribological (lateral force) properties of amorphous/nanocrystalline solid materials [27]. But, investigations are concentrated separately either on nanoindentation or on nanoscratch of coatings and therefore, a conclusive correlation between mechanical and tribological properties of the Fe-based coatings is lacking.
ACCEPTED MANUSCRIPT In the present work, Fe-based amorphous powder with lean composition was used to deposit amorphous/nanocrystalline composite coating on mild steel substrate through atmospheric plasma spraying. The effect of spraying power i.e. the heat input on the microstructure, distribution of different nanocrystalline secondary phases in the amorphous matrix, and its
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effect on the mechanical and tribological properties of the coatings have been correlated
systematically. An analytical model developed by Chen et al. [27] has been used to correlate the mechanical and tribological properties of the coatings. The aim of the present paper is to
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provide theoretical support for defining the suitable spraying power in depositing Fe-based
properties.
2. Experimental procedures 2.1. Coating deposition
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amorphous/nanocrystalline composite coating with better mechanical (hardness and wear)
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Gas atomized Fe-based amorphous powders (Fe-2.5Cr-6.7Si-2.5B-0.7C, wt%) procured from Epson Atmix Corporation, Japan was used as the feedstock for coating deposition. This chosen alloy composition possesses high potential for industrial applications because of its
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high glass forming ability and relatively low cost of the alloying elements. The coatings with a thickness of about 150 ± 40 µm were deposited by an atmospheric plasma spraying system
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(AMTAG Adv. Mat. Tech., Switzerland, MP 200 Pulse Arc 1000) equipped with a 6-axes robotic arm. The feedstock powders were sprayed onto cold rolled close annealed (CRCA) mild steel substrates of dimension 50 mm x 20 mm x 2 mm. Before spraying process, the substrates were grit-blasted using 24 grit silica beads and ultrasonically cleaned with ethyl alcohol. Argon (99.9 % pure) and Hydrogen (99.9 % pure) of constant flow rates were used as primary gas and secondary gas, respectively; whereas the plasma power was varied in the range of 18.6–35.5 kW by keeping plasma current fixed and changing voltage during the
ACCEPTED MANUSCRIPT plasma spraying process. Thus, the plasma power values were obtained as 18.6 kW, 23.7 kW, 29 kW and 35.5 kW while actually aiming for 20 kW, 25 kW, 30 kW and 35 kW, respectively. Higher plasma power than 35 kW was not tried in this work, as higher plasma power leads to high degree of oxidation which promotes devitrification. Besides, smaller
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feedstock powders (mostly in the range of 10 to 30 µm were selected for this work to get the advantage of better inter-splat bonding resulting from uniform melting of lower size powders possessing higher specific surface area. However, lower particles also experience higher in-
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flight temperature which promotes the self-annealing effect after deposition, which could lead to even more devitrification. Therefore, the plasma power was not raised beyond 35 kW,
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keeping in mind the effect of (i) degree of oxidation, (ii) uniform melting and (iii) selfannealing in the deposited coatings. The details of the plasma spraying parameters are shown in Table 1. For convenience, coatings prepared under different plasma spraying conditions are
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denoted as Coating-I, Coating-II, Coating-III, and Coating-IV, respectively.
2.2. Microstructural characterization
The surface morphology and microstructure of the coatings were investigated with a scanning
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electron microscope (SEM, SUPRA 40, Carl Zeiss AG, Germany). The porosity of the
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coatings was evaluated by quantitative image analysis performed on the SEM images using ImageJ k 1.45 software, where area analysis method was applied by calculating the ratio of porosity area to the total area of a particular field of view (FOV). The primary factor to precisely differentiate the gray scale of porosities form that of the rest of the microstructural features is conducting the process of image thresholding meticulously. At first, the primitive SEM images were converted into 32-bit gray scale images. Then, FFT-based band pass filtering was done before performing thresholding operation to minimize the noise which may get introduced into the image near the porosity boundaries after thresholding. A baseline
ACCEPTED MANUSCRIPT thresholding operation was carried out with a carefully selected threshold parameters based on the pixel value histogram [28]. At least ten SEM images with a magnification of 300X were obtained from the polished cross-sections surface of the various coatings and were used subsequently as FOVs for the porosity estimation. The phase formation in the feedstock
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powders and coatings was characterized by X-Ray diffraction (XRD, PANalytical-Empyrean) equipped with a monochromatic Cr-Kα (λ= 0.229 nm) X-ray source in the range of 30° ≤ 2θ ≤ 110°. The phase analysis obtained from XRD results were validated by carrying out
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transmission electron microscopy (FEG-TEM, JEOL, JEM-2100F) on the coatings. For TEM observation, samples of about 0.4 mm thickness were machined out by wire-cut machine
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from the as-sprayed coating surface, perpendicular to the coating deposition direction. These EDM cut samples were thinned down to 30 µm by mechanical grinding, followed by ionbeam milling using a Gatan (model 691) system at an incident angle of 4-6° with a beam
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energy of 4 KeV until perforation.
2.3. Microindentaion and nanoindentation experiments Vickers micro indenter (UHLVMHT - 001, Walter Uhl, Germany) was used with a 25 gf load
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and a dwell time of 10 s to evaluate the microhardness of the sintered samples.
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Microhardness values were obtained from the average of 15 individual measurements performed on the polished top surface of the coatings. Nanoindentation was carried out using a Triboindenter (TI 950, Hysitron Inc., USA) at room temperature with a standard three-sided pyramid Berkovich tip (TI-0039, Hysitron Inc., total included an angle of 142.3º and tip radius of 100 nm) at a peak load of 5000 µN and loading-unloading rate of 500 µN/s. Depthsensing indentation test was carried out on a standard fused silica (Poisson's ratio of 0.17) specimen beforehand for the calibration of the indenter tip area function. The average nanohardness and reduced modulus reported in this paper represent the average of 100
ACCEPTED MANUSCRIPT indents (10 X 10 matrix). The indenter was held at peak load for 10 s to remove any creep effect. The spacing between indentations was set to be 30 µm to avoid the plastic flow effect of previous indentations on subsequent indentations.
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2.4. Nanoscratch and dry sliding wear tests
Nanoscratch test was conducted under ramp loading with a peak force of 5 N at a loading rate of 125 mN s-1 and indenter velocity of 5 µm s-1 over a scratch length of 200 µm. Dry sliding
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wear test on the coatings was conducted by using a pin-on-disc tribometer (TE97 Friction and Wear Demonstrator, Phoenix Tribology Ltd., England) to evaluate the tribological properties.
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The coatings as pin were held against the counter face of a rotating alumina disc with wear track diameter 25 mm. Before tests, the coatings were sequentially ground with 600, 1000, 1200, 1500, and 2000 mesh SiC papers. The polished coatings were ultrasonically cleaned in acetone for 5 min and then dried in air. The sliding tests were conducted at room temperature
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with a sliding distance of 235 m, a load of 10 N, and a constant speed of 0.26 m s−1. The wear rate was calculated by using equation, W = V/SF, where W is the wear rate in mm3/Nm, V is the wear volume loss in mm3, S is the total sliding distance in m, and F is the applied load in
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N. At least three repeat scratch tests were carried out for each coating sample under
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predefined test conditions.
3. Results and discussions 3.1. Morphology and phase structure of feedstock powders Fig. 1(a) shows the SEM micrograph of the Fe-based gas-atomized powders used as feedstock during plasma spraying in the present study. Most of the powder particles are of spherical morphology and exhibit smooth surface, which conforms to the morphological features of powders synthesized by gas atomization. It can be observed from the bar chart of
ACCEPTED MANUSCRIPT particle size distribution, shown in Fig. 1(b) that the particle size varies in the size range of 5 to 55 µm, however most of the particles are of 10-30 µm. Spherical morphology of powders with a smooth surface and this particular size range are indicative of good flowability during the thermal spraying process. It was also found that some of the larger particles were capped
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with smaller ones, as shown by arrows in Fig. 1(a). When particles with different sizes collide to each other in the gas turbulence during atomization process, then smaller particles having higher solidification rate compare to the larger one, would tend to adhere to the molten
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surface of the larger particles [29]. Fig. 1(c) presents the XRD pattern of the feedstock
powders. Presence of a broad halo peak appearing in 2θ range of 50° - 90° without any
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crystalline peak indicates that the powders were amorphous in nature. The formation of amorphous phase is attributed to the high GFA of the Fe−2.5Cr−6.7Si−2.5B−0.7C (wt.%) alloy and fairly high solidification rate during gas atomization of the powders in an inert
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atmosphere.
3.2. Microstructure of as-sprayed coatings
Top surface SEM images of the as-sprayed coatings deposited with varying plasma power
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(18.6, 23.7, 29 and 35.5 kW) are presented in Fig. 2(a-d) to show the splat morphology after
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deposition of the coatings; whereas the polished surface images in Fig. 2(e-h) depict the internal microstructure. It can be clearly seen from Fig. 2(a) that Coating-I, which was deposited at a low spraying power (18.6 kW, Table 1), contained higher extent of partially molten particles on the unpolished coating surface. The corresponding image (Fig. 2e) of the polished surface exhibits high amount of large pores. As the spraying power was gradually increased from 18.6 kW to 35.5 kW, the amount of partially molten particles reduced (Fig. 2b-d). Higher degree of powder melting in case of maximum plasma power of 35.5 kW is evident from the corresponding Fig. 2d. It can also be observed from the micrographs of the
ACCEPTED MANUSCRIPT polished coating surface (Fig. 2e-h) that the amount and size of porosity were gradually reduced as the plasma power increased. The cross-sectional morphology of the various coatings is shown in Fig. 3. All of the four coatings with thickness of around 150 ± 40 µm were well-bonded to the mild steel substrate and the inter-splat bonding increased gradually
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from Coating-I to Coating-IV (Fig. 3a-d), as the plasma power increased. The porosity of coatings tends to decline from Coating-I to Coating-IV and the porosity vol.% were 14.8%, 12.6%, 9.1%, and 4.8% for Coating-I, Coating-II, Coating-III, and Coating-IV, respectively
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(as reported in Table 2). The low spraying power was inadequate to completely melt the
feedstock powders and therefore led to deposition of poorly adhered splats which resulted in
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higher porosity in Coating-I (Fig. 3a). The Coating-ІV possessed the lowest because of welladhered splats (Fig. 3d). The dense structure of this coating is mainly ascribed to the higher degree of melting leading to better spreading of splats upon impact on the substrate and previously deposited splat layers [13, 21, 30]. Higher spraying power provides more enthalpy
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to accelerate the particle melt and reduces the surface tension of droplets; and therefore, the droplets flatten well, leading to the formation of a dense coating structure [19]. Besides, at higher spraying power, gas expands fast and leads to high speed plasma flow and hence
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drives the molten particles to impact intensely on substrate and to form denser coating [21].
3.3. Phase composition of the coatings Fig. 4 shows XRD patterns of the different coatings obtained at different spraying powers. XRD patterns of all the four coatings exhibit a broad halo peak appearing in 2θ range of 55– 85° with some low-intensity crystalline peaks. Thus, the XRD results insinuate synthesis of amorphous/nanocrystalline composite coatings. Comparing the XRD patterns of the feedstock powders (Fig. 1b) and the coatings (Fig. 4), it is evident that precipitation of crystalline phases occurred during the coating deposition process. The crystalline peaks
ACCEPTED MANUSCRIPT appeared in the XRD patterns of the coatings are associated with α-Fe (PDF No. #00-0060696), Fe23B6 (PDF No. #00-047-1332) and B13C2 (PDF No. #00-026-0233) phases. The reasons behind crystalline phase formation can be attributed to the following factors. Firstly, depending on the heat input and associated degree of melting caused by increase in plasma
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power, certain amount of crystallization took place during splat solidification. Secondly, a certain amount of heat build-up took place due to deposition of successive molten splats and release of latent heat of fusion during solidification of these splats [19, 21, 30]. Besides, there
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was a probability of powder oxidation to some extent during the coating deposition attributed to high-temperature processing in atmospheric environment which may lead to the faster
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crystallization during solidification [21, 26]. An et al. also reported increased crystallization in Fe-based metallic glass system with higher degree of oxidation [21]. The amorphous phase content of the samples was estimated from the area ratio of crystalline peaks to amorphous hump [31, 32]. Based on this calculation, retained amorphous phase fraction in Coating-I,
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Coating-II, Coating-III, and Coating-IV were 86 %, 83.6%, 82.8%, and 81%, respectively. This indicates that the fraction of amorphous phase in coatings decreased with increase in plasma arc power. Although there is a trend of increasing crystallinity with increasing plasma
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power, however, the variation is not very significant. The higher amount of heat input in case
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of higher plasma power deposition would lead to more degree of melting (evident from Fig. 2), higher amount of heat build-up in the deposited splats and would have higher probability of oxide formation. All these factors contributed in higher amount of crystallinity in coatings deposited with higher plasma power. It must be mentioned here that phase amount estimation only from XRD patterns is not enough to ensure if a material is completely amorphous. Usually, the identification of small amounts of nanocrystalline phases embedded in high amorphous content matrix could not be clearly identified from the XRD measurements due to the weak intensity peaks of these phases. Moreover, these weak intensity peaks can be
ACCEPTED MANUSCRIPT confused with the background and/or be overlapped by the amorphous phase broad halo peak. The authors intend to carry out exhaustive differential scanning calorimetry on the coatings to find out the amount of retained amorphous phase, where a ribbon of similar composition synthesized by splat melting at a very high wheel speed would be used as the standard fully
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amorphous material.
TEM was employed to understand the phase evolution within the coatings at different plasma power and to validate the composite structure (presence of nanocrystalline and amorphous
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phases) of the coatings. TEM micrographs and corresponding selected area diffraction (SAD) patterns of Coating-I, Coating-II, and Coating-IV are shown in Fig. 5. The TEM micrographs
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(Fig. 5a-c) of these coatings exhibit presence of nanocrystalline phases uniformly distributed in the amorphous matrix and this observation matches well with the corresponding SAD patterns (Fig. 5d-f), which show spotted rings (obtained from nanocrystalline phases) with a diffused background (from amorphous matrix). The TEM images (Fig. 5a-c) reveal that grain
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number density and grain size in the coating increased when the plasma power increased from 18.6 kW to 35.5 kW. This finding confirms that volume fraction of various nanocrystalline phases increased with increasing plasma power and this is consistent with the
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corresponding XRD patterns (Fig.4) showing gradual decrease in hump intensity with
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increasing plasma power.
The HRTEM image obtained from Coating-IV shown in Fig. 6a clearly exhibits the presence of nanocrystalline precipitates embedded in the amorphous matrix which further confirms the composite nature of the produced coating. The selected areas (circles 1 to 4) have been digitally analyzed to obtain the fast Fourier transform (FFT) patterns and reported in Fig. 6(be). Formation of various crystalline phases was envisaged from these FFT patterns. The diffraction spots in Fig. 6(b and c) fit to the α-Fe (1 1 0) plane (d= 2.02 Å) and (2 0 0) plane (d= 1.43 Å) forming an angle of 45° along the [0 0 -1] zone axis, and cubic Fe23B6 (2 4 2)
ACCEPTED MANUSCRIPT plane (d= 2.19 Å) and (4 4 0) plane (d= 1.90 Å) forming an angle of 30° along the [1 -1 1] zone axis, respectively. These d-spacings are well in accord with the d-spacings of various phases which were identified from the XRD patterns presented earlier in Fig. 4. Formation of additional phases, as shown in Fig. 6(d and e), which are tetragonal Fe2B (0 2 0) plane (d=
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2.55 Å) and (1 2 1) plane (d= 2.01 Å) forming an angle of 35° along the [1 0 -1] zone axis and orthorhombic Fe3B showing (1 0 2) plane (d= 2.03 Å) and (2 2 0) plane (d= 2.10 Å) forming an angle of 72° along the [-2 2 1] zone axis, respectively. Formation of similar
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intermetallic phases in Fe-Si-B amorphous systems during devitrification has been reported
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earlier by other researchers [33-36].
3.4. Microindentaion and nanoindentation results of coatings
The mechanical properties including Vickers microhardness, nanohardness, and reduced elastic modulus of the coatings and the mild steel substrate obtained from microindentaion
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and nanoindentation tests are listed in Table 2. All of the coatings exhibited higher microhardness as well as nanohardness than the substrate. The Vickers microhardness of substrate, Coating-I, Coating-II, Coating-III, and Coating-IV were 172 ± 10 HV0.025, 799 ±
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124 HV0.025, 895 ± 123 HV0.025, 954 ± 110 HV0.025, and 993 ± 86 HV0.025, respectively.
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Moreover, the nanohardness of substrate, Coating-I, Coating-II, Coating-III, and Coating-IV were 3.61 ± 0.25 GPa, 9.60 ± 2.30 GPa, 10.51 ± 2.42 GPa, 11.00 ± 2.52 GPa, and 11.80 ± 2.71 GPa, respectively, and the reduced elastic modulus were 184.9 ± 5.6 GPa, 125.6 ± 34.3 GPa, 138.5 ± 38.8 GPa, 154.9 ± 30.8 GPa, and 167.9 ± 31.5 GPa, respectively. It is evident from this data that there is an increase in hardness (both micro and nano) and reduced elastic modulus of the coatings with increase in spraying power. The higher hardness and modulus observed in coatings deposited with higher plasma power can be attributed to (i) the better inter-splat bonding resulting lower porosity of coatings, as shown in Fig. 3 and Table 2, and
ACCEPTED MANUSCRIPT (ii) homogeneously precipitated hard intermetallic phases (Fe2B, Fe3B, and Fe23B6) in the amorphous matrix, (as apparent from the XRD - Fig. 4 and HRTEM - Fig. 6 analysis) which resulted in dispersion strengthening. However, average elastic modulus of the Fe-based metallic glass coatings (125.6 to 167.9 GPa) is considerably lower than mild steel substrate
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(184.9 GPa). This may be due to crystalline (α-Fe) nature of the mild steel substrate and
highly amorphous structure of the coatings. Amorphous alloys have large free volume which increases the inter-atomic distance and weakens the atomic bonding. After crystallization,
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free volume annihilates and results in a relaxed structure with shorter inter-atomic distance, leading to higher hardness and Young's modulus [33, 36-39]. Sun et al. [39] observed that
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partially or fully crystallized structure was showing higher hardness and Young's modulus than the material with a completely amorphous structure. Similar behavior has been observed in present work in terms of hardness and modulus. Average nanohardness increased from 9.6 GPa to 11.8 GPa, and reduced elastic modulus increased 125.6 GPa to 167.9 GPa as
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crystalline content increased at elevated plasma power. Lashgari et al [36] reported that Febased fully amorphous ribbon had reduced elastic modulus of around 160 GPa and increased up to 240 GPa after annealing, due to precipitation of Fe2B and Fe23B6 compounds which are
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extremely brittle. Average reduced elastic modulus of the coatings in the present work
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corroborates well with the earlier reported values [36, 39]. For the coatings having multiphase structure, each individual phase affects the overall properties of the coatings. In this context, grid indentation is a very useful technique for characterization of materials of small volume and to provide quantitative prediction of phase transformations. This two-dimensional surface mapping technique with large arrays of nanoindentation is useful to investigate the properties of individual constituent phases by carrying out statistical analysis of the resulting data (nanohardness and elastic modulus). Resulting nanohardness can be plotted in the form of a histogram of indent number density
ACCEPTED MANUSCRIPT with different hardness range which is associated with different phases present in that material. Based on the nanohardness value reported in previous literatures [13, 36, 39-41], the range of nanohardness data was categorized into four sets viz. 5.0-7.5 GPa, 7.5-9.5 GPa, 9.5-12.5 GPa, and 12.5-20.0 GPa, belonging to (i) mixed nano-sized crystalline α-Fe and Fe-
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based amorphous matrix, (ii) amorphous matrix, (iii) mixed amorphous and nano-sized Febased intermetallic phases, and (iv) nano-sized intermetallics rich regions, respectively. It can be noted here that, the Fe-based amorphous phase has been considered in the first three sets,
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as a wide range of hardness for Fe-based amorphous phase based on presence of various
alloying elements has been reported in various literatures. Recently, Lashgari et al. reported
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the presence of similar nanohardness range for amorphous phase and mixed amorphous and nano-sized borides in Fe-Si-B based alloy system synthesized by melt spinning [36]. Graphical representation of indent number density (%), as shown in Fig. 7, shows the relative fraction of different phases present in the various coatings. From Fig. 7, it can be observed
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that the indent number density (%) with hardness range of 7.5-9.5 GPa, which represents the hardness of amorphous matrix kept on decreasing with increase in plasma power. On the other hand, this percentage for the nano-sized intermetallics rich region with hardness range
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of 12.5-20.0 GPa showed an increasing trend with increasing heat input. It can be commented
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that at elevated spraying power, indenter falling on mixed phase (amorphous and nano-sized intermetallic phases) and nano-sized intermetallics phase has higher probability due to more fraction of nano-sized intermetallics formation. In case of Coating-I, 17% and 30% of nanohardness values were in the range of 5.0-7.5 (mixed nano-sized α-Fe and amorphous matrix) and 7.5-9.5 GPa (amorphous matrix), respectively, in comparison to 5.4% and 9.8% in Coating-IV. The percentage of nanohardness data (9.5-12.5 GPa) related to the mixed amorphous and intermetallic phase in Coating-I and Coating-IV (45.4% and 48.9%, respectively) remained almost similar. Moreover, coating-ІV exhibited more amount of
ACCEPTED MANUSCRIPT intermetallic phases in comparison to Coating-I as 35.9% of hardness data was in 12.5-20.0 hardness range, in comparison to 7.8% for Coating-I. The variation in nanohardness values can also be justified from the TEM analysis. Further, presence of the different phases can be validated by the typical load-displacement curves obtained from grid nanoindentation test as
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plotted in Fig. 8(a). The depth of penetration and corresponding nanohardness values for different type of phases has been depicted in Fig. 8(b). Moreover, maximum indentation
depth was approximately 188 nm for mixed α-Fe and amorphous matrix phase region, and
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approximately 106 nm for the intermetallics region at a constant peak force of 5000 µN, as shown in Fig. 8(b). The decrease in the penetration depth could be justified by the fact that
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intermetallic regions show higher nanohardness and stiffness (increase in the slope of unloading portion of the load-displacement curve). Fig. 8(c) shows 3D depth profile images obtained by in-situ scanning probe microscopy (SPM) carried out on four different regions with indentations on the coating surface. The four different penetration depths (127 nm, 110
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nm, 88 nm and 67 nm) represent the four different hardness range regions (mixed nano-sized α-Fe and amorphous matrix, amorphous phase, mixed amorphous and intermetallic phases, and intermetallic phases) as mentioned earlier. Therefore, the variation in indentation depth
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(from 67 to 127 nm, approximately) validates the presence of several phases in the coating.
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Besides hardness and Young’s modulus, nanoindentation can also be used to predict the wear resistance, which is related to the ratio of H/Er of the coatings [33, 42-44]. The values of H/Er ratio obtained from the nanoindentation with a load of 5000 µN are listed in Table 2. It has been found that coatings have higher H/Er (H/Er ~ 0.07) value than the steel substrate (H/Er ~ 0.019). Many authors have suggested that higher value of the ratio H/Er, generally indicative of the good wear resistance [13, 33, 42-44]. Mukhopadhyay et al. reported that H/Er ratio more than 0.05 indicates high degree of elastic recovery in the material which is another important parameter for higher wear resistance [44]. From the obtained H/Er values of the
ACCEPTED MANUSCRIPT coatings, it appears that coatings are suitable for the applications where high degree of elastic recovery is required and it should exhibit better wear resistance than the mild steel substrate. The trend of H/Er values will not only depend on the individual H and Er values, but also on
trend of H/Er values became inconsistent.
3.5. Nanoscratch behavior of the coatings
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the rate of change of H and Er. As the rate of increase in H and Er values were not similar, the
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Fig. 9 shows SEM images of the worn-out surfaces of the various coating after the nano-
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scratch test. The worn surfaces of the coatings indicated that grooves became wider with increasing load in scratch direction and coatings deposited at various spraying power have the different capability to resist the splats detachment. Coating-I and Coating-II which were deposited at lower plasma arc powers, displayed severe splat detachments (Fig. 9a and b), due to the lower degree of melting of the feedstock powders and associated poor inter-splat
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bonding. However, due to higher degree of melting of feedstock powders at higher spraying powers (28.6 and 35.5 kW), Coating-III and Coating-IV (Fig. 9c and d) experienced
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improved inter-splats bonding and thus better scratch resistance with no or a very little splats detachment in comparison with Coating-I and Coating-II. Thus, it could be concluded that
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high spraying power resulted in more stable scratch which resulted in comparatively smoother scratch profile (Fig. 9c and d). Additionally, fracture of coating surface is brittle in nature and therefore it can be concluded that coatings are liable to brittle fracture during the friction process. Presence of various nano-sized intermetallics and their distribution in the amorphous matrix would play important role in dictating the wear behavior of the coatings and this issue should be addressed in future work in this area. To understand the tribological behavior of the coatings, dry sliding wear test was carried out, the results of which are reported later in Section 3.7.
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3.6. Correlation between nanoindentation and nanoscratch results Mechanical properties (viz. nanohardness and reduced elastic modulus) and tribological properties (lateral force and scratch path morphology) of the plasma sprayed coatings were
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greatly affected by the change in the spraying power, as shown by the results reported in
Table 2 and Fig. 9. Therefore, correlation between mechanical and tribological properties of the coatings was required to be investigated. Recently, Chen et al. developed an analytical
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model to correlate nano-scale mechanical properties with tribological properties of coatings,
2
1/ 2 Fmax H τ F f = 24.5 hm − 2 2.263 × Er
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which is given in equation (1) [27]. (1)
Where, Ff is lateral force (tribological property) obtained during the scratch test, H is hardness, Er is reduced elastic modulus, and τ is shear strength. The effect of surface
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roughness and material sinking-in/pile up on the real contact area during loading has not been included in the above-proposed model. Based on the lateral force values obtained from the
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nanoscratch tests, and elastic modulus and hardness obtained from nanoindentation tests, the shear strength (τ) of the mild steel substrate and the different coatings was estimated using
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Equation 1. The shear strength was found to be 188 ± 12 MPa for the mild steel substrate, 398 ± 56 MPa for Coating-I, 433 ± 46 MPa for Coating-II, 516 ± 54 MPa for Coating-III, and 670 ± 60 MPa for Coating-IV, as listed in Table 2. Coating-IV has significantly high shear strength than other coatings and approximately 3.6 times greater than the mild steel substrate. It can be concluded that higher scratch resistance of the coatings resulted due to the higher shear strength of the respective coatings. Higher shear strength of the coatings leads to higher lateral force during scratch of the coatings, which finally results in higher scratch resistance.
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3.7. Tribological behavior of coatings accessed by dry sliding wear test Fig. 10 depicts the results obtained from pin-on-disc based dry sliding wear tests carried out on the coatings and mild steel substrate at a load of 10 N and sliding velocity of 0.26 m/s.
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Fig. 10(a) illustrates the variation in coefficient of friction (COF) of coatings and substrate as a function of sliding distance. The COF values of the mild steel substrate, Coating-I, CoatingII, Coating-III, and Coating-IV were 0.47 ± 0.06, 0.26 ± 0.05, 0.20 ± 0.02, 0.18 ± 0.03, and
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0.15 ± 0.02, respectively, as reported in Table 2. All coatings showed lower COF than that of the substrate, owing to the higher hardness of the coatings as reported in Section 3.4.
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Moreover, the average COF of the coatings slightly decreased with increasing plasma power as mentioned above. Relatively higher fluctuations in COF curves and higher value of COF were observed for Coating-I compared to other coatings. This was mainly due to its poorly bonded splats (Fig. 3a) with relatively bigger pores (Fig. 2e). The lowest coefficient of
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friction for Coating-IV can be attributed to the high hardness which resulted from the tightly stacked splats (Fig. 3d) and precipitation of higher fraction of very hard Fe-borides phases, as was shown earlier in Fig. 8.
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Fig. 10(b) shows the wear rate of the mild steel substrate and the various coatings. The wear rate of the substrate, Coating-I, Coating-II, Coating-III, and Coating-IV were 37.4 ± 0.9, 16.1
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± 2.2, 10.2 ± 1.5, 8.2 ± 0.8, and 3.9 ± 1.6 mm3/Nm, respectively. The substrate showed the maximum wear rate compared to that of the coatings. A decreasing trend of wear rate of the coatings was observed with increasing plasma arc power. Moreover, the wear rate of the Coating-IV (3.9 mm3/Nm) decreased approximately one-tenth of that of the substrate (37.4 ± 0.9 mm3/Nm), which confirms that the wear resistance of Coating-IV was much superior to that of the substrate. Higher value of H/Er ratio, as discussed in Section 3.5 also supported the fact of higher wear resistance of the coatings. Thus, the deposition of Fe-based
ACCEPTED MANUSCRIPT amorphous/nanocrystalline composite coating deposited at 35.5 kW plasma power (CoatingIV) could effectively protect the mild steel substrate from friction wear and therefore prolong its service life.
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4. Conclusions
Fe-based amorphous/nanocrystalline composite coatings were developed on a mild steel substrate by atmospheric plasma spraying (APS) of gas-atomized amorphous powders
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(Fe−2.5Cr−6.7Si−2.5B−0.7C, wt%) with varying plasma power. The coatings primarily
consisted of amorphous phase, which was due to the high glass forming ability (GFA) of the
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selected multi-component system and the high cooling rate of APS technique. Higher plasma power resulted in higher degree of melting causing better inter-splat bonding and precipitation of higher fraction of nanocrystalline intermetallics (Fe23B6, Fe2B, and Fe3B) in the amorphous matrix. The nanoscratch results indicated that the increased plasma power
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resulted in uniform scratch profiles due to well-bonded splats. Moreover, dry sliding wear test showed that both coefficient of friction and wear rate of the coatings decreased with increasing plasma power. In the present study, the highest nanohardness of 11.8 ± 2.71 GPa
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and the minimum wear rate of 3.9 ± 1.6 mm3/Nm was achieved for the coating deposited at
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the plasma power of 35.5 kW. An analytical model which was used to correlate the mechanical and tribological properties of the coatings predicted that the coating with higher plasma power has the higher shear strength which ultimately resulted in higher scratch resistance. The present results show that Fe-based amorphous/nanocrystalline composite coatings have great potential for the industrial applications in the areas that involve high wear.
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Acknowledgements The author, T. Laha thankfully acknowledges the financial support obtained from Research
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and Development Division of Tata Steel, India.
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Coating-ІІ
Coating-ІІІ 60
Secondary gas (H2) flow rate (L/min)
10
Powder feed rate (g/min)
30
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Primary gas (Ar) flow rate (L/min)
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Carrier gas (Ar) flow rate (L/min) Current (A) Voltage (V)
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200
300
400
500
93
79
72.5
71
18.6
23.7
29
35.5
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Plasma power (kW)
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Coating thickness (µm)
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Coating-I
Coating-II
Coating-III
Coating-IV
Porosity (vol.%)
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14.8 ± 2.6
12.6 ± 1.7
9.1 ± 0.8
4.8 ± 0.6
Vickers microhardness
172 ± 10
799 ± 124
895 ± 123
954 ± 110
993 ± 86
Nanohardness (GPa)
3.61 ± 0.25
9.60 ± 2.30
10.00 ± 2.42
11.00 ± 2.52
Reduced elastic modulus (GPa)
184.9 ± 5.6
125.6 ± 34.3
148.3 ± 38.8
149.5 ± 30.8
167.9 ± 31.5
Wear property parameter, H/Er
0.019 ± 0.001
0.076 ± 0.025
0.067 ± 0.033
0.074 ± 0.049
0.071 ± 0.010
Shear strength (MPa)
188 ± 12
398 ± 56
433 ± 46
516 ± 54
670 ± 60
Coefficient of friction (Dry Sliding wear)
0.47 ± 0.06
0.26 ± 0.05
0.20 ± 0.02
0.18 ± 0.03
0.15 ± 0.02
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Fig. 2. SEM micrographs of as-sprayed (a-d) and polished (e-h) top surface morphology of the plasma sprayed coatings deposited with varying plasma power (18.6, 23.7, 29 and 35.5 kW, showing the variation in degree of melting and associated porosity in the various
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Fig. 3. SEM images of cross-sections of as-sprayed iron-based amorphous/nanocrystalline
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coatings deposited with plasma power of (a) 18.6 kW, (b) 23.7 kW, (c) 29 kW and (d) 35.5 kW, showing the variation in morphology.
Fig. 4. XRD patterns of the various plasma sprayed coatings deposited with varying plasma power showing variation in crystallization.
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Fig. 5. TEM images of Coating-I, Coating-II, Coating-ІIІ, and Coating-ІV (a - c) and the respective SAED patterns (d - f).
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Fig. 6. High-resolution transmission electron microscopy (HRTEM) results of Coating-IV: (a) HRTEM image confirming the uniform distribution of nano-sized grains embedded in
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amorphous matrix, (b-e) FFT patterns obtained from the different locations (viz. 1, 2, 3 and 4) in Fig. a corresponding to various crystalline phases i.e. α-Fe, Fe23B6, Fe3B, and Fe2B phases, respectively.
Fig. 7. Bar graph showing the indent number density (%) of the plasma sprayed coatings in four different nanohardness range, viz. 5.0-7.5 GPa, 7.5-9.5 GPa, 9.5-12.5 GPa and 12.5-20.0 GPa, which belong to mixed nano-sized α-Fe and amorphous matrix, amorphous matrix,
ACCEPTED MANUSCRIPT mixed amorphous and nano-sized intermetallics and nano-sized intermetallics rich regions, respectively. Fig. 8. (a) Load vs. penetration depth plots obtained from nanoindentation test corresponding to different phases (amorphous and nano-sized crystalline) present in the Fe-based metallic
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glass composite coatings, (b) the depth of penetration and corresponding nanohardness
obtained from the different phases of Coating-IV deposited at 35.5 kW plasma power and (c) 3D depth profile images obtained by in-situ scanning probe microscopy (SPM) carried out on
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four different indentation regions.
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Fig. 9. SEM images of the scratched surface of the different coatings after ramping load nano-scratch tests at 5 N peak load and 5 µm/s scratch velocity: (a) Coating-І, (b) Coating-ІI, (c) Coating-ІIІ, and (d) Coating-ІV, respectively.
Fig. 10. (a) Coefficient of friction (COF) plotted as a function of sliding distance and (b)
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wear rate obtained from pin-on-disc based dry sliding wear tests carried out on the coatings
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and mild steel substrate at a load of 10 N and sliding velocity of 0.26 m/s.
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ACCEPTED MANUSCRIPT Highlights Fe-based amorphous/nanocrystalline coatings were deposited by plasma spraying.
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Coatings consist of nano-sized α-Fe and Fe-borides distributed in amorphous matrix.
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Higher plasma power resulted in more devitrification and better intersplat bonding.
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High hardness resulted due to denser coating and presence of hard Fe-borides.
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Wear resistance of the coatings increased with increasing plasma power.
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•