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Microstructure and oxidation behavior of NiCr-chromium carbides coating prepared by powder-fed laser cladding on titanium aluminide substrate
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Microstructure and oxidation behavior of NiCr-chromium carbides coating prepared by powder-fed laser cladding on titanium aluminide substrate S.E. Aghilia,b,∗, M. Shamaniana, R. Amini Najafabadic, A. Keshavarzkermanib, R. Esmaeilizadehb, U. Alib, E. Marzbanradb, E. Toyserkanib a
Department of Materials Engineering, Isfahan University of Technology, 84156-83111, Isfahan, Iran Multi-Scale Additive Manufacturing Lab, Department of Mechanical and Mechatronics Engineering, University of Waterloo, Waterloo, Ontario, N2L 3G1, Canada c Department of Metallurgy and Materials Engineering, Golpayegan University of Technology, Golpayegan, Iran b
A R T I C LE I N FO
A B S T R A C T
Keywords: Powder-fed laser cladding Additive manufacturing Oxidation NiCr–Cr3C2 Titanium aluminide Oxidation mechanism
In the present study, a NiCr–Cr3C2 powder mixture was prepared by mechanical alloying and then coated on titanium aluminide substrates by the powder-fed laser cladding process using a set of optimum parameters. The high temperature oxidation behavior of the substrate and coating was studied by isothermal annealing at 900 °C for 5 h. It was found that the microstructure of the coating is composed of γ solid solution with different chromium carbide phases (Cr3C2, Cr7C3 and Cr23C6). The presence of different chromium carbides in the microstructure of coating can be attributed to the partial melting of primary Cr3C2 and the formation of nonequilibrium carbide phases during rapid cooling of laser cladding. The NiCr-chromium carbide laser cladded coating samples showed superior oxidation resistance compared to the substrate. The oxidation mechanism of both coating and substrate follow the parabolic law, where the parabolic rate constant of the coating was 20% of that of the substrate at 900 °C. Time-of-Flight Secondary Ion Mass Spectroscopy (ToF-SIMS) and Grazing Angle X-Ray Diffraction (GAXRD) analysis revealed that the surface of the oxide layer formed on the NiCr-chromium carbides coating and the substrate is mostly composed of Cr2O3 and TiO2, respectively.
1. Introduction Titanium aluminides are well known candidates to replace nickel based super alloys in high temperature applications due to the low density while providing desirable high temperature strengths [1,2]. However, low ductility at room temperature and unsuitable wear and oxidation resistance at high temperatures are several drawbacks, which may limit their applications in various circumstances. It is reported that titanium aluminides containing 40–48 wt% Al show a dual phase microstructure composed of TiAl and Ti3Al that exhibit better ductility at room temperature [3]. Heat treatment and thermo-mechanical processes during the production of titanium aluminides are other techniques, which are used to refine the microstructure and improve the mechanical properties at room temperature [4]. Furthermore, titanium aluminides demonstrate poor oxidation and wear resistance at high temperatures. Their oxidation resistance drops drastically at 800 °C due to the formation of porous TiO2 layers with low adhesion to the substrate [5]. The wear rate of titanium aluminides is also high due to its relatively low hardness (∼320 HV). Several studies have been conducted to improve the high ∗
temperature properties of titanium aluminides. Surface treatment techniques can be considered as one of the most functional methods to overcome the aforementioned problems due to the simplicity and effectiveness. Amongst all possible surface treatment options, the application of laser cladding is growing fast [6,7]. Laser cladding, while still an expensive process due to maintenance and initial costs, has the advantage of providing a strong metallurgical bonding between the coating and the substrate with minimum dilution [6]. In addition, the overall heat input in laser cladding process is low compared to the other techniques resulting to a thin heat affected zone [6]. The porosity of laser cladded coatings are also low compared to thermal spray techniques which is essential for oxidation resistance. All these advantages have made this process an effective platform for additive manufacturing as well, when the process is called laser directed energy deposition (LDED). Many efforts have been presented to manufacture effective coatings on TiAl substrate using the laser cladding process [8–10]. The wear performance of laser cladded TiAl/TiC coatings on TiAl substrate was studied by Wang et al. [8]. Their results showed that the hardness of the cladding increased to 756 HV after the optimization of process
Corresponding author. Department of Materials Engineering, Isfahan University of Technology, 84156-83111, Isfahan, Iran. E-mail addresses: [email protected], [email protected] (S.E. Aghili).
https://doi.org/10.1016/j.ceramint.2019.09.139 Received 19 August 2019; Received in revised form 11 September 2019; Accepted 15 September 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: S.E. Aghili, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.09.139
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2.2. Microstructure analysis
parameters, compared to 430 HV of the substrate. Liu et al. [9] investigated the wear and oxidation behavior of Ti5Si3/γ/TiSi composite cladding on TiAl substrate and reported a 15% improvement in the wear resistance after cladding. The effect of Co–Cr–Mo coating on the oxidation behavior of TiAl was also investigated by Barekat et al. [10]. A dense duplex layer consisting of Cr2O3 and CoCr2O4 was found to form on the surface of cladding after exposure to oxidizing environment at 900 °C for prolonged duration. Recently, Cr3C2 and WC based coatings have been introduced as an alternative for the hard-chromium coatings as the conventional hard chromium coatings are not temperature resistant and are also carcinogenic due to their Cr(VI) compounds [11]. Cr3C2– NiCr coated components can be operated up to 900 °C, which is much higher than the safe working temperature of WC-Co cermets (∼450 °C) [12]. Pan et al. [13] studied the effect of chromium carbide content (35 and 50 wt.%) on the mechanical properties of Cr3C2–Ni60 laser cladded coating and reported higher hardness for the coating containing 50 wt% Cr3C2. Yang et al. [14] has also reported improved room and high temperature wear behavior of steel substrate after laser cladding of NiCr-C-WS2 composite powder. Although a few examinations have been performed to improve the high temperature properties of TiAl substrate by laser cladding of carbide based coatings, the literature lacks knowledge about the microstructure and high temperature oxidation behavior of the powder-fed laser cladded NiCr–Cr3C2 process. In the present study, NiCr–Cr3C2 powder mixture was synthesized by mechanical alloying and subsequently cladded on the surface of titanium aluminide (TiAl) substrate using the powder-fed laser cladding process when an optimum set of parameters was used. The microstructure and oxidation behavior of coating was carefully examined. Results show that NiCr-chromium carbides offer a superior oxidation resistance than titanium aluminide.
X-Ray diffraction (XRD-Philips XPERT MPD) was used to identify the phases during MA and laser cladding process. For Grazing angle XRD (GAXRD) analysis, the ω was considered as 0.6° to generate the Xray from the near surface regions. The microstructure of samples was characterized using a field emission scanning electron microscope (FESEM; MIRA3, TESCAN, Czech Republic) and a Transmission electron microscope (TEM-120 kV Philips CM120 and JEOL 2010F). A mixture of HF–HNO3–H2O with the volume ratio of 1-6-7 was used as etching reagent. Electron energy loss spectroscopy (EELS) was employed for quantitative elemental analysis in TEM. The grain structure and phase maps of the optimal clad microstructure were also performed by electron backscattered diffraction, EBSD (JEOL 7000F SEM with an Oxford detector). The surface roughness of the laser cladded coatings were measured by a confocal microscope (Keyence VK-X250) with three repetitions on an area of 700 × 500 μm. X-ray computed tomography (CT) (ZEISS Xradia 520 Versa) was employed to study the distribution and size of the porosities in the cladded layer and overlapped claddings. The final 3D image was constructed using a beam hardening constant of 0.05 and analyzed in Dragonfly 3.1 (Object Research Systems (ORS), Montreal, Canada) software. Nanoindentation tests were performed using Wrexhan, UK equipped with Berkovich diamond indenter under load control mode. A total of 20 indentations were carried out for each sample to measure the average hardness and the elastic modulus. 2.3. Oxidation tests The high temperature oxidation behavior of TiAl substrate and coating was examined by Thermo-gravimetric analysis (TGANETZSCH, STA 449 F1 Jupiter) at 900 °C for 5 h in stationary air with a heating rate of 50 °C/min on 3✕3✕0.3 mm samples. Samples were carefully cleaned with acetone before oxidation test to remove the effect of oil/contaminations on the weight change during the test. TGA was repeated three times on each sample and the average readings were reported. The samples were characterized by SEM and XRD after oxidation test to understand the acting mechanisms during high temperature oxidation. TOF-SIMS (time-of-flight secondary ion mass spectroscopy) technique was also used to study the oxidation mechanisms through analyzing the type of negative ions on the oxidized surface. Analysis was performed using spectrometry mode with Bi3+ ions used for analysis and Cs (1 keV) ions used for sputtering on an analysis area of 350 μm2 and sputter area of 150 μm2.
2. Materials and methods 2.1. Preparation of substrate and coatings In this work, titanium aluminide was used as the substrate. The chemical composition of the substrate is presented in Table 1 and comprises of Ti and Al with small additions of Cr and Nb. Fig. 1 shows the XRD pattern of the titanium aluminide substrate. The substrate is composed of a mixture of TiAl and Ti3Al phases. Surface pretreatment of substrate was performed by grinding the substrate with SiC papers and cleaning with acetone to remove the surface oxides and/or contaminations before laser cladding. NiCr–Cr3C2 powder mixture was synthesized by mechanical alloying (MA) and rounded by spray drying in order to obtain spherical particles due to better flow properties [15]. The coaxial laser cladding process was performed in an argon environment using a powder feed system with a 0.7 KW transverse-flow fiber laser system equipped with a 4-axes CNC machine. Argon has been used as the shielding and powder carrier gas with the flow rate of 150 mL/min. For all experiments, the laser-beam spot size has been fixed of 2 mm. A constant laser power of 400 W with different scanning speeds (2–5 mm/s) and powder feeding rates (200–400 mg/s) were used to clad the as-milled NiCr–Cr3C2 powder mixture on the TiAl substrate to form single tracks. The final overlay for the proposed coating was achieved using a side-by-side overlapping (50%) of single track cladding with the optimal condition.
3. Results and discussion 3.1. Synthesis and characterization of NiCr–Cr3C2 powder mixture by mechanical alloying As mentioned before, Ni and Cr elemental powders were milled up to 40 h to prepare Ni–Cr solid solution. Fig. 2a shows the XRD patterns of Ni-20 wt% Cr powder mixture in as-mixed condition and after different milling times up to 40 h. The XRD pattern of as-mixed powder mixture consisted of Ni and Cr elemental peaks. Enhancement of the milling time to 10 h resulted in the reduction of peaks intensity and an increase in the peak width. This could be related to the crystallite size reduction and internal stress increment of the as-received elements [16–18]. Furthermore, Ni peaks shift slightly to the lower diffraction angles with increasing the milling time from 5 h to 30 h, confirming the dissolution of Cr atoms in the Ni crystal structure and formation of Ni (Cr) solid solution over time. The position of Ni peaks remain constant with further milling after 30 h indicating that further milling could only reduce the crystallite size of Ni(Cr) solid solution. The crystallite size of the nanostructured Ni(Cr) after 30h of milling was predicted as 43 nm using Williamson-Hall equation [19]. TEM analysis was also performed
Table 1 Chemical composition (wt.%) of the titanium aluminide substrate. Element
Ti
Al
Cr
Nb
Others
Chemical composition (wt.%)
60.54
32.51
2.27
4.65
0.03
2
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Fig. 1. XRD pattern of titanium aluminide substrate.
with Cr3C2. Chromium and graphite powders were also mixed with the molar ratio of 3:2 and milled up to 40 h in order to synthesize the Cr3C2 compound. Fig. 3 shows the XRD patterns of as-received and milled Cr–C powder mixture. Reduction of the intensity of chromium and graphite peaks are observed during mechanical alloying for up to 7 h.
to observe the microstructure of the formed Ni(Cr) compound after 30 h milling time (Fig. 2c). Nanometric grains are clearly visible in the individual powder particles shown in Fig. 2c. More information about the synthesis mechanism of the Ni(Cr) solid solution can be found elsewhere [20]. Since the dissolution of Cr in Ni is completed after 30 h of MA, the 30 h milled Ni(Cr) powder was used for subsequent mixture
Fig. 2. (a) XRD patterns of Ni-20 wt% Cr powder mixture at different milling times up to 40 h, (b) Magnified X-ray diffraction (XRD) peak around 2θ = 50–54, (c) TEM micrograph of the nanostructured Ni(Cr) solid solution after 30 h. 3
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Fig. 3. XRD patterns of Cr–C powder mixture with the molar ratio of 3:2 at different milling times.
powder mixture. The NiCr-50 wt.% Cr3C2 powder mixture was then rounded with spray drying to obtain spherical particles with the average size of around 50–120 μm, resulting in better flow during powder fed laser cladding.
By increasing the milling time to 10 h, a Cr3C2 peak appears beside the Cr peaks. Interestingly, further milling resulted to the formation of Cr7C3 instead of Cr3C2. Therefore, the XRD pattern of 40 h milled chromium-graphite powder mixture shows only Cr7C3 diffraction peaks. The same behavior (formation of Cr7C3 and Cr23C6 instead of Cr3C2) is also observed by Gomari et al. [21] during the milling of chromium-graphite powder mixture with the molar ratio of 3:2. Thermodynamic analysis was performed to understand the origin of such phenomenon. According to Equations (1) and (2) [21], the formation of both Cr3C2 and Cr7C3 chromium carbide compounds from the elemental powders is possible during mechanical alloying due to negative Gibbs free energy of formation. It is worth mentioning that the Gibbs free energy change of Cr7C3 at the milling temperature is more negative than that of Cr3C2. Graphite is a very ductile material with high tendency to stick to the milling media and balls during mechanical alloying. This leads to small changes of the molar ratio of Cr: C, resulting in the formation of Cr7C3 instead of the desirable Cr3C2 compound.
3Cr + 2C → Cr3 C2 ΔG (Cr3 C2) = −83189 − 0.967T J /mol ΔH298 = −72.3 kJ / mol
(1)
7Cr + 3C → Cr7 C3 ΔG (Cr7 C3) = −142666 − 30.92T J /mol ΔH 298 = −144.4 kJ / mol
(2)
3.2. Characterization of laser cladded coatings Fig. 5 shows the SEM micrographs of the single track cross-section of laser cladded coatings produced at laser power of 400 W with different scanning speed (S) and powder feeding rate (F). At a first glance, it can be seen that the geometrical properties of single track laser cladded coatings are highly dependent on the process parameters. No successful claddings were produced at higher powder feeding rate and low scanning speed due to detachment from the substrate (Fig. 5). Meanwhile, the powder-fed laser cladding process for all other combinations of scanning speed and powder feeding rate resulted in the formation of a sound cladding with unique geometrical properties. It should be mentioned that the overall heat input increases with the reduction of scanning speed, resulting in a high instability in the melt pool that is detrimental for the formation of high quality clad. Fig. 5 shows a decrease in clad height with increasing scanning speed and decreasing powder feeding rate. At the constant power of 400 W, the width of claddings was decreased with increasing the scanning speed. The reason can be explained by the reduction of the interaction time between the powder particles and laser beam at higher scanning speed [22]. In this case, the generated heat input to the centre of the molten pool decreases and as a consequence, the width of the claddings decreases, too. The enhancement of the powder feeding rate can also absorb some of the laser energy and results to lower heat input and therefore smaller clad width [22]. It can be observed that some of the coatings contain numerous porosities in their cross-section. Considering geometrical requirements and porosity content of the coatings, the sample cladded with the scanning speed of 3 mm/s and powder feeding rate of 300 mg/s was selected as the optimal cladding and considered for further characterization. It is worth to mention that the optimization of the laser cladding parameters for NiCr–Cr3C2 on TiAl substrate is comprehensively described in Ref. [22]. The surface of the optimal single track was studied to measure the surface roughness (Ra) (Fig. 6a-b). Roughness is an important cladding parameter which can influence the fatigue properties of the coating [23,24]. The fatigue mechanism can be divided into three stages: crack
As mentioned before, the milled Cr–C powder mixture is composed of the elemental peaks up to 7 h of milling time. Therefore, a two-step procedure consisting of 7 h milling and 3 h annealing at 900 °C was employed in order to fabricate Cr3C2. Fig. 4a shows the XRD patterns of 7 h milled Cr–C powder mixture after subsequent annealing. Cr3C2 diffraction peaks observed in Fig. 4a clearly show the presence of the more desirable Cr3C2 compound due to the reaction of activated chromium and graphite during annealing. Fig. 4b shows the TEM micrograph of the synthesized Cr3C2 after annealing. EELS analysis was employed in order to quantify the concentrations of Cr and C in the formed compound. It was found that the synthesized chromium carbide after annealing is composed of ∼40 at.% C which matches with the nominal composition of Cr3C2. The synthesized Ni(Cr) and Cr3C2 powder mixtures were mixed with equal weight ratio and further milled for 1 h to obtain a homogenous 4
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Fig. 4. (a) XRD patterns of Cr–C powder mixture after 7 h of milling and subsequent annealing at 900 °C for 3 h (b) TEM micrograph of the synthesized Cr3C2 after annealing.
initiation, cracks propagation and final failure. The first stage, crack initiation, plays an important role in low stress fatigue samples. It is well known that the rough surfaces normally exhibit poor fatigue properties due to the role of surface defects on the cracks initiation [23,24]. The average surface roughness of the single track surface was measured as 3 μm which seems to be sufficient for most practical applications. X-ray computed tomography (CT) was also performed on the coating and shows that the optimal cladding contains only 0.2% porosity with a median pore size of less than 20 μm (Fig. 6c). The microstructure of different parts of the single track cladding
was carefully studied by EBSD. Fig. 7 shows the SEM micrograph and EBSD result of the interface of titanium aluminide (TiAl) substrate and NiCr-chromium carbides single track coating. A thin layer of Ti3Al with the average thickness of ∼8 μm is formed between the substrate and cladding and acts as the interface. The microstructure of the coating close to the interface is very complex and is composed of different phases (Fig. 7b). Presence of different chromium carbide phases (Cr23C6 and Cr7C3) are observed close to the interface. In addition, the presence of TiC, Ti8C5 and Al4C3 close to the interface is probably caused by the dilution effect of the substrate. As the coating is applied by laser
Fig. 5. SEM micrographs of the cross section of the single-track laser cladded NiCr–Cr3C2 powder mixture with the laser power of 400 W and different scanning speed (S) and powder feeding rate (F). 5
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Fig. 6. (a) Surface profile of the optimal NiCr-chromium carbides single track cladding before and after curvature corrections (a–b) along with the result of computed tomography (CT) scan (c).
cladding, the surface layers of the substrate melt and elements from the substrate dilute in the cladded layer. The microstructure of the optimal cladding far away from the interface was also studied which shows a mixture of γ solid solution Cr7C3, Cr23C6 and the un-melted Cr3C2 particles having an average size of ∼10 μm (Fig. 8). Although, the as-prepared powder mixture has the chemical composition of NiCr-50 wt% Cr3C2, different carbide phases were detected on the cross-section of optimal cladding. One possibility for such behavior might arise from the partial melting of Cr3C2 during laser cladding and the formation of non-equilibrium chromium carbides upon solidification. The formation of non-equilibrium phases during the laser cladding of Cr3C2–NiCr feedstock powder have been also reported by Yang et al. [14]. Complete transformation of Cr3C2 to Cr7C3 was reported by Yang et al. [14]. Heubner et al. [25] has also reported partial dissolution of WC fine particles during laser cladding of IN625WC composite powder, resulted to the formation of different secondary carbides with higher amounts of Nb, Cr and W. It is worth mentioning that, in this investigation, some traces of un-melted Cr3C2 was observed mostly on the regions close to the top surface of the cladding which is the last region of solidification. In fact, the Cr3C2 particles are partially melted and solidified to Cr7C3 and Cr23C6 and the remaining un-melted particles floated on the melt and remain there. It should be noted that TiC, Ti8C5 and Al4C3 were not observed away from the substrate. Nanoindentation tests were performed to assess the effectiveness and performance of the cladded layer. Fig. 9 shows the load-displacement nanoindentation curve of the substrate and the optimal cladding. Since, the nanoindentation test was conducted in load control mode, lower penetration depth of the cladding reflects its higher hardness compared to the titanium aluminide substrate. The hardness and elastic
Fig. 7. (a) SEM micrograph of the interface of substrate and the optimal single track coating (b) EBSD phase map of the area close to the interface. 6
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H 2 ⎛1 E ln ⎜ tanα ⎞⎟ =K+ σy 3 ⎝ 3 σy ⎠
(3)
where H and E represent the hardness and elastic modulus obtained from nanoindentation test, σy is the yield strength, α is a constant angle defined as 19.7° for Berkovich indenter and K = 4/3. Other values of K by Studman and Bushby [28], were reported as 7/6 and 1.15. Fig. 9 shows the yield strength of the substrate and the optimal laser cladded coating based on different cavity models [27,28]. The value of yield strength is in the same range for all models, about 1.1 GPa for TiAl and 3.6 GPa for the optimal cladding. Based on the results from nanoindentation tests and yield strength, it can be concluded that the mechanical behavior of the substrate is improved by laser cladding. Since the final cladded surface layer forms from the overlapping of single tracks, its microstructure was carefully studied to correlate it to the mechanical properties oxidation resistance. Fig. 10a shows the cross-section of the overlapped NiCr-chromium carbides coating on titanium aluminide substrate. After overlapping, the average thickness of the coating increases to ∼500 μm. SEM analysis reveals presence of voids in the overlapped coating. A detailed CT analysis of the overlapped coating showed a 2.3% porosity and is presented in Fig. 10b. CT results show a higher porosity than the optimal cladded layer but a lower porosity than thermal spray deposition techniques such as Air plasma spray (APS) [29] and high-velocity oxy fuel (HVOF) [30,31]. Repeated melting of the coating during overlapping may cause turbulence in the melt pool, resulting in a higher amount of trapped gasses during solidification [32,33]. It was observed that most of the porosities have a radius less than 40 μm and only limited number of pores have larger radii. The XRD pattern of the overlapped coating is shown in Fig. 11 and confirms that the coating is composed of a mixture of γ and different chromium carbides (Cr3C2, Cr7C3 and Cr23C6). 3.3. Evaluation of oxidation resistance
Fig. 8. (a) SEM micrograph, (b) EBSD phase map of the NiCr-chromium carbides single-track cladding at the middle height of coating.
Fig. 12 shows the specific mass gain (mass gain per unit area) versus oxidation time for the substrate and overlapped coating at 900 °C. As seen, the specific mass gain of the coating is lower than the substrate. The specific mass gain of the laser cladded coating is ∼0.004 mg/cm2, twice lower than ∼0.009 mg/cm2 for the substrate. It should be noted that, for both samples, the slope of the mass gain curve is always higher at the beginning of oxidation and reduces with increasing oxidation time. To quantify the rate of oxidation, the curve of (mass gain per unit area)2 versus oxidation time was plotted for both substrate and coating and is presented in Fig. 12b. As can be observed, a straight line fits the data points from TGA experiments, meaning that the oxidation kinetics is parabolic. The value of R2, for the experimental data of substrate and
modulus of NiCr-chromium carbides cladding were recorded as 9.5 ± 1.2 GPa and 250 ± 11 GPa, respectively which is higher than that of TiAl substrate (hardness = 3.6 ± 0.2 GPa and elastic modulus = 164 ± 4 GPa). The presence of different types of chromium carbides in the microstructure of the optimal cladding is responsible to the remarkable increase of the hardness of the coating, due to the higher hardness of carbide particles compared to the Ni and/or NiCr alloys [26]. Johnson cavity model [27] was used to estimate the yield strength of the substrate and cladding from nanoindentation data (Equation (3)).
Fig. 9. Typical load-displacement curve of TiAl substrate and NiCr-Chromium carbides optimal cladding during nanoindentation test under load-control mode and yield strength (GPa) of the TiAl substrate and NiCr-chromium carbide single track cladding estimated by cavity models. 7
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Fig. 10. (a) Cross section overview of NiCr-chromium carbides overlapped coating produced by laser cladding and (b) its related computed tomography (CT) scan showing the porosity size.
compared to the substrate. It is reported that Cr2O3 is stable up to the temperature of 950 °C, and therefore can protect the surface coating at high temperatures [37]. There are some minor peaks related to different carbides and also γ solid solution which might be a result of the original coating. It is reported that the thickness of oxide film on the surface of Cr3C2-25 wt%NiCr is less than 2 μm after 16 days of oxidation at 800 °C and do not form a continuous surface layer before 8 h of oxidation [38]. Therefore, the peaks related to γ and different chromium carbide phases from the XRD pattern of oxide layer might be generated due to the effect of original coating. The formation of Cr23C6 carbide phase is also previously reported [12] during isothermal oxidation of Cr3C2-25% NiCr at 800 °C and attributed to the stepwise oxidation of Cr3C2. The SEM morphology of the coating surface after oxidation test also confirms that the formed layer is not continuous and some parts of the original coating are still visible on the surface. SIMS studies were carried out to study the oxidation mechanisms of the substrate and coating. Fig. 15a–b shows the SIMS depth profiles of titanium aluminide and NiCr-chromium carbides oxidized at 900 °C for 5 h, respectively. Results from the titanium aluminide substrate show that the intensity of TiO− and TiO2− is high at the surface of oxide, while decreasing continuously across the depth and instead the intensity of AlO− and AlO2− negative ions is grown. It further confirms that the oxidation of titanium aluminide starts with the formation of Al2O3 and then continues with the formation of TiO2 which is in good agreement with the results of the XRD and EDS measurements on the top surface of the oxide layer. It is worth to mention that the crystal lattice of TiO2 contains both oxygen vacancies and interstitial titanium atoms. The former prevails at low temperatures and high oxygen pressure and the latter at high temperatures and low oxygen pressure. It is reported that at T < 800 °C, TiAl is mainly oxidized through diffusion of oxygen from the oxidizing media towards the interface of substrate/oxide layer [35]. However, the oxygen partial pressure of inner
coating was calculated to be 0.94 and 0.985, respectively. A parabolic oxidation rate can be represented by Equation (4), where “t” is oxidation time and K and C are the kinetics constants [34]. (Δm/a)2 = kpt + C
(4)
The slope of fitted straight lines (parabolic rate constant-Kp) for substrate and coating is then calculated and given in Fig. 12b. The parabolic rate constant of the coating is 5 times smaller than the substrate indicating the improved oxidation resistance of substrate after laser cladding. The morphology of oxide layer formed on the surface of titanium aluminide after 5 h of oxidation is shown in Fig. 13(a–b). A dense layer of oxidation products was formed on the surface of substrate after oxidation. The GAXRD pattern reveals that the outer surface of the oxide layer is mostly composed of TiO2 with some minor amount of Al2O3 (Fig. 13c). EDS analysis on the surface of oxide layer also shows that the oxide layer is composed of ∼60 at.% O, 36.5 at.% Ti and 2.5 at. % Al with small additions of Nb and Cr. These results confirm that the outer surface of the oxide layer is almost fully covered with TiO2 (Rutile). It is reported that the affinity of Ti and Al to oxygen is very close [35]. In the other words, both TiO2 and Al2O3 have similar thermodynamic stability. However, the growth rate of TiO2 is faster than Al2O3 due to its unique crystal structure which provides extra sites for diffusion of oxygen. However, the oxide layer is composed of both Al2O3 and TiO2 at the early stages of oxidation, TiO2 nuclei grow faster and the outermost layer of oxide layer is enriched with rutile [36,37]. The morphology and XRD pattern of oxide layer formed on the surface of coating after oxidation test is shown in Fig. 14. Results in Fig. 14 show that chromium oxide (Cr2O3) is the major phase in the oxide layer. The formation of chromium oxide-rich scale in the present study probably acts as a barrier to the diffusion of oxygen into the coating and therefore improved oxidation resistance of coating
Fig. 11. XRD pattern of NiCr-chromium carbides overlapped coating. 8
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Fig. 12. (a) Mass gain per unit area and (b) (mass gain per unit area)2 versus oxidation time for TiAl substrate and NiCr-chromium carbides coating during TG experiments at 900 °C and parabolic rate constant of substrate and coating in high temperature oxidation test at 900 °C.
2 The microstructure of the coating was mainly composed of γ and different chromium carbide phases (Cr3C2, Cr7C3 and Cr23C6). The formation of different chromium carbide phases might arise from the partial melting of Cr3C2 powders and formation of non-equilibrium Cr7C3 and Cr23C6 phases. 3 The NiCr-chromium carbides coating showed a superior oxidation resistance at 900 °C compared to the titanium aluminide substrate. 4 The oxidation behavior of NiCr-chromium carbides coating and titanium aluminide substrate follow the parabolic law. The parabolic rate constant of the coating was one fifth of the substrate at 900 °C. 5 ToF-SIMS analysis revealed that the oxidation of substrate started with the formation of Al2O3 and further continued with growth of TiO2. In contrast, the surface of oxide layer on the NiCr-chromium carbide coating was almost fully covered with Cr2O3.
layers decreases with increasing the temperature and/or thickness of the oxide layer, resulting to the high concentration of titanium interstitials which promotes further diffusion of Ti towards outer layers [35]. In the case of NiCr-chromium carbide coating (Fig. 15b), it can be observed that the intensity of CrO−, CrO2− and CrO3− is very high at the surface and seems to be constant throughout the thickness of the oxide layer. It is reported in literature that the CrO− ions provide the best clue for the presence of Cr2O3 in the oxidized region [39]. Therefore, it can be concluded that Cr2O3 is the main compound in the NiCr-chromium carbide coating. 4. Conclusions In the present study, the NiCr–Cr3C2 powder mixture was synthesized by mechanical alloying and then coated on the TiAl substrate using the powder-fed laser cladding process. The following conclusions can be drawn from this study:
Acknowledgements The authors would like to thank Canadian Centre for Electron Microscopy (CCEM) and Department of Mechanical Engineering at McMaster University for conducting the EBSD, TEM and Nanoindentation measurements. The authors gratefully thank WATLAB at the University of Waterloo for conducting the SIMS experiments. The authors would also like to thank the members of Multi-Scale Additive
1 The optimized cladding was obtained at P = 400 W, S = 3 mm/s and F = 300 mg/s that resulted in single tracks with approximately 0.2% porosity. The final overlay for oxidation test was achieved through the side-by side overlapping (50%) of the single track cladding with the optimal laser condition. 9
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Fig. 13. (a–b) Morphology and (c) GAXRD pattern of oxide layer formed on the surface of titanium aluminide after high temperature oxidation test at 900 °C for 5 h.
Fig. 14. Morphology and GAXRD pattern of oxide layer formed on the surface of NiCr-chromium carbides coating after high temperature oxidation test at 900 °C for 5 h. 10
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Fig. 15. SIMS profile of the (a) titanium aluminide substrate and (b) NiCr-chromium carbide coating after high temperature oxidation test.
Manufacturing (MSAM) group especially Prof. Mihaela Vlasea, Dr. Vladimir Paserin, Dr. Hamed Asgari and Dr. Yahya Mahmoodkhani for their supports.
[7] J. Kusinski, S. Kac, A. Kopia, A. Radziszewska, M. Rozmus-Gornikowska, B. Major, L. Marczak, A. Lisieck, Laser modification of the materials surface layer – a review paper, Bull. Pol. Acad. Sci. Chem. 60 (2012) 4. [8] J. Wang, R.G. Song, X. Lin, W.D. Huang, Microstructure and properties of laser cladding TiC/TiAl composite coatings on gamma-TiAl intermetallic alloy, Surf. Eng. 25 (2009) 196–200. [9] X.B. Liu, H.M. Wang, Microstructure, wear and high-temperature oxidation resistance of laser clad Ti5Si3/γ/TiSi composite coatings on γ-TiAl intermetallic alloy, Surf. Coat. Technol. 200 (2006) 4462–4470. [10] M. Barekat, R. Shoja Razavi, A. Ghasemi, High temperature oxidation behavior of laser clad Co–Cr–Mo coating on γ-TiAl substrate, J. Laser Appl. 28 (2016) 042005. [11] J.M. Guilemany, N. Espallargas, P.H. Suegama, A.V. Benedetti, Comparative study of Cr3C2–NiCr coatings obtained by HVOF and hard chromium coatings, Corros. Sci. 48 (2006) 2998–3013. [12] H. Wang, X. Yan, H. Zhang, M. Gee, Ch. Zhao, X. Liu, X. Song, Oxidation-dominated wear behaviors of carbide-based cermets: a comparison between WC-WB-Co and Cr3C2-NiCr coatings, Ceram. Int., in press. https://doi.org/10.1016/j.ceramint. 2019.07.113. [13] Ch. Pan, H. Wang, H. Wang, Q. Chang, H. Wang, Microstructure and thermal physical parameters of Ni60-Cr3C2 composite coating by laser cladding, J. Wuhan Univ. Technol.-Materials Sci. Ed. 25 (2010) 991–995. [14] M. Sh. Yang, X.B. Liu, J.W. Fan, X.M. He, Sh.H. Shi, G.Y. Fu, M.D. Wang, Sh.F. Chen, Microstructure and wear behaviors of laser clad NiCr/Cr3C2–WS2 high temperature self-lubricating wear-resistant composite coating, Appl. Surf. Sci. 258 (2012) 3757–3762. [15] S. Liu, Y. Zhang, R. Kovacevic, Numerical simulation and experimental study of powder flow distribution in high power direct diode laser cladding process, Lasers Manuf. Mater. Process. 2 (2015) 199–218. [16] B.D. Cullity, Elements of X-Ray Diffraction, second ed., Addison-Wesley, Menlo
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[29] V. Sreenivasulu, M. Manikandan, High-temperature corrosion behaviour of air plasma sprayed Cr3C2-25NiCr and NiCrMoNb powder coating on alloy 80A at 900 °C, Surf. Coat. Technol. 337 (2018) 250–259. [30] G. Bolelli, L.-M. Berger, T. Börner, H. Koivuluoto, V. Matikainen, L. Lusvarghi, C. Lyphout, N. Markocsan, P. Nylén, P. Sassatelli, R. Trache, P. Vuoristo, Sliding and abrasive wear behaviour of HVOF- and HVAF-Sprayed Cr3C2–NiCr hardmetal coatings, Wear vols. 358–359, (2016) 32–50. [31] S. Matthews, M. Hyland, B. James, Microhardness variation in relation to carbide development in heat treated Cr3C2–NiCr thermal spray coatings, Acta Mater. 51 (2003) 4267–4277. [32] Y. Zhao, T. Yu, Ch. Guan, J. Sun, X. Tan, Microstructure and friction coefficient of ceramic (TiC, TiN and B4C) reinforced Ni-based coating by laser cladding, Ceram. Int. 45 (2019) 20824–20836. [33] M. Erfanmanesh, H. Abdollah-Pour, H. Mohammadian-Semnani, R. ShojaRazavi, Kinetics and oxidation behavior of laser clad WC-Co and Ni/WC-Co coatings, Ceram. Int. 44 (2018) 12805–12814. [34] S. Mrowec, A. Stoklosa, Calculations of parabolic rate constants for metal oxidation, Oxid. Metals 8 (1974) 379–391. [35] V.G. Chuprina, I.M. Shalya, Oxidation of TiAl intermetallics, Powder Metall. Met. Ceram. 46 (2007) 11–12. [36] L. Zhao, Y. Zhang, Y. Zhu, J. Lin, Z. Ren, Oxidation behavior of γ-TiAl based alloys at 900°C in air, Mater. High. Temp. 33 (2016) 234–240. [37] P. Elliott, A practical guide to high-temperature alloys, Mater. Des. 12 (1991) 299–307. [38] Y. Ding, T. Hussain, D.G. McCartney, High-temperature oxidation of HVOF thermally sprayed NiCr–Cr3C2 coatings: microstructure and kinetics, J. Mater. Sci. 50 (2015) 6808–6821. [39] X. Wu, S. Voyshnis, A. Seyeux, Y. Chumlyakov, Ph. Marcus, ToF-SIMS study of oxide films thermally grown on nickel-base alloys, Corros. Sci. 141 (2018) 175–181.
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Microstructure and oxidation behavior of NiCr-chromium carbides coating prepared by powder-fed laser cladding on titanium aluminide substrate S.E. Aghilia,b,∗, M. Shamaniana, R. Amini Najafabadic, A. Keshavarzkermanib, R. Esmaeilizadehb, U. Alib, E. Marzbanradb, E. Toyserkanib a
Department of Materials Engineering, Isfahan University of Technology, 84156-83111, Isfahan, Iran Multi-Scale Additive Manufacturing Lab, Department of Mechanical and Mechatronics Engineering, University of Waterloo, Waterloo, Ontario, N2L 3G1, Canada c Department of Metallurgy and Materials Engineering, Golpayegan University of Technology, Golpayegan, Iran b
A R T I C LE I N FO
A B S T R A C T
Keywords: Powder-fed laser cladding Additive manufacturing Oxidation NiCr–Cr3C2 Titanium aluminide Oxidation mechanism
In the present study, a NiCr–Cr3C2 powder mixture was prepared by mechanical alloying and then coated on titanium aluminide substrates by the powder-fed laser cladding process using a set of optimum parameters. The high temperature oxidation behavior of the substrate and coating was studied by isothermal annealing at 900 °C for 5 h. It was found that the microstructure of the coating is composed of γ solid solution with different chromium carbide phases (Cr3C2, Cr7C3 and Cr23C6). The presence of different chromium carbides in the microstructure of coating can be attributed to the partial melting of primary Cr3C2 and the formation of nonequilibrium carbide phases during rapid cooling of laser cladding. The NiCr-chromium carbide laser cladded coating samples showed superior oxidation resistance compared to the substrate. The oxidation mechanism of both coating and substrate follow the parabolic law, where the parabolic rate constant of the coating was 20% of that of the substrate at 900 °C. Time-of-Flight Secondary Ion Mass Spectroscopy (ToF-SIMS) and Grazing Angle X-Ray Diffraction (GAXRD) analysis revealed that the surface of the oxide layer formed on the NiCr-chromium carbides coating and the substrate is mostly composed of Cr2O3 and TiO2, respectively.
1. Introduction Titanium aluminides are well known candidates to replace nickel based super alloys in high temperature applications due to the low density while providing desirable high temperature strengths [1,2]. However, low ductility at room temperature and unsuitable wear and oxidation resistance at high temperatures are several drawbacks, which may limit their applications in various circumstances. It is reported that titanium aluminides containing 40–48 wt% Al show a dual phase microstructure composed of TiAl and Ti3Al that exhibit better ductility at room temperature [3]. Heat treatment and thermo-mechanical processes during the production of titanium aluminides are other techniques, which are used to refine the microstructure and improve the mechanical properties at room temperature [4]. Furthermore, titanium aluminides demonstrate poor oxidation and wear resistance at high temperatures. Their oxidation resistance drops drastically at 800 °C due to the formation of porous TiO2 layers with low adhesion to the substrate [5]. The wear rate of titanium aluminides is also high due to its relatively low hardness (∼320 HV). Several studies have been conducted to improve the high ∗
temperature properties of titanium aluminides. Surface treatment techniques can be considered as one of the most functional methods to overcome the aforementioned problems due to the simplicity and effectiveness. Amongst all possible surface treatment options, the application of laser cladding is growing fast [6,7]. Laser cladding, while still an expensive process due to maintenance and initial costs, has the advantage of providing a strong metallurgical bonding between the coating and the substrate with minimum dilution [6]. In addition, the overall heat input in laser cladding process is low compared to the other techniques resulting to a thin heat affected zone [6]. The porosity of laser cladded coatings are also low compared to thermal spray techniques which is essential for oxidation resistance. All these advantages have made this process an effective platform for additive manufacturing as well, when the process is called laser directed energy deposition (LDED). Many efforts have been presented to manufacture effective coatings on TiAl substrate using the laser cladding process [8–10]. The wear performance of laser cladded TiAl/TiC coatings on TiAl substrate was studied by Wang et al. [8]. Their results showed that the hardness of the cladding increased to 756 HV after the optimization of process
Corresponding author. Department of Materials Engineering, Isfahan University of Technology, 84156-83111, Isfahan, Iran. E-mail addresses: [email protected], [email protected] (S.E. Aghili).
https://doi.org/10.1016/j.ceramint.2019.09.139 Received 19 August 2019; Received in revised form 11 September 2019; Accepted 15 September 2019 0272-8842/ © 2019 Elsevier Ltd and Techna Group S.r.l. All rights reserved.
Please cite this article as: S.E. Aghili, et al., Ceramics International, https://doi.org/10.1016/j.ceramint.2019.09.139
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2.2. Microstructure analysis
parameters, compared to 430 HV of the substrate. Liu et al. [9] investigated the wear and oxidation behavior of Ti5Si3/γ/TiSi composite cladding on TiAl substrate and reported a 15% improvement in the wear resistance after cladding. The effect of Co–Cr–Mo coating on the oxidation behavior of TiAl was also investigated by Barekat et al. [10]. A dense duplex layer consisting of Cr2O3 and CoCr2O4 was found to form on the surface of cladding after exposure to oxidizing environment at 900 °C for prolonged duration. Recently, Cr3C2 and WC based coatings have been introduced as an alternative for the hard-chromium coatings as the conventional hard chromium coatings are not temperature resistant and are also carcinogenic due to their Cr(VI) compounds [11]. Cr3C2– NiCr coated components can be operated up to 900 °C, which is much higher than the safe working temperature of WC-Co cermets (∼450 °C) [12]. Pan et al. [13] studied the effect of chromium carbide content (35 and 50 wt.%) on the mechanical properties of Cr3C2–Ni60 laser cladded coating and reported higher hardness for the coating containing 50 wt% Cr3C2. Yang et al. [14] has also reported improved room and high temperature wear behavior of steel substrate after laser cladding of NiCr-C-WS2 composite powder. Although a few examinations have been performed to improve the high temperature properties of TiAl substrate by laser cladding of carbide based coatings, the literature lacks knowledge about the microstructure and high temperature oxidation behavior of the powder-fed laser cladded NiCr–Cr3C2 process. In the present study, NiCr–Cr3C2 powder mixture was synthesized by mechanical alloying and subsequently cladded on the surface of titanium aluminide (TiAl) substrate using the powder-fed laser cladding process when an optimum set of parameters was used. The microstructure and oxidation behavior of coating was carefully examined. Results show that NiCr-chromium carbides offer a superior oxidation resistance than titanium aluminide.
X-Ray diffraction (XRD-Philips XPERT MPD) was used to identify the phases during MA and laser cladding process. For Grazing angle XRD (GAXRD) analysis, the ω was considered as 0.6° to generate the Xray from the near surface regions. The microstructure of samples was characterized using a field emission scanning electron microscope (FESEM; MIRA3, TESCAN, Czech Republic) and a Transmission electron microscope (TEM-120 kV Philips CM120 and JEOL 2010F). A mixture of HF–HNO3–H2O with the volume ratio of 1-6-7 was used as etching reagent. Electron energy loss spectroscopy (EELS) was employed for quantitative elemental analysis in TEM. The grain structure and phase maps of the optimal clad microstructure were also performed by electron backscattered diffraction, EBSD (JEOL 7000F SEM with an Oxford detector). The surface roughness of the laser cladded coatings were measured by a confocal microscope (Keyence VK-X250) with three repetitions on an area of 700 × 500 μm. X-ray computed tomography (CT) (ZEISS Xradia 520 Versa) was employed to study the distribution and size of the porosities in the cladded layer and overlapped claddings. The final 3D image was constructed using a beam hardening constant of 0.05 and analyzed in Dragonfly 3.1 (Object Research Systems (ORS), Montreal, Canada) software. Nanoindentation tests were performed using Wrexhan, UK equipped with Berkovich diamond indenter under load control mode. A total of 20 indentations were carried out for each sample to measure the average hardness and the elastic modulus. 2.3. Oxidation tests The high temperature oxidation behavior of TiAl substrate and coating was examined by Thermo-gravimetric analysis (TGANETZSCH, STA 449 F1 Jupiter) at 900 °C for 5 h in stationary air with a heating rate of 50 °C/min on 3✕3✕0.3 mm samples. Samples were carefully cleaned with acetone before oxidation test to remove the effect of oil/contaminations on the weight change during the test. TGA was repeated three times on each sample and the average readings were reported. The samples were characterized by SEM and XRD after oxidation test to understand the acting mechanisms during high temperature oxidation. TOF-SIMS (time-of-flight secondary ion mass spectroscopy) technique was also used to study the oxidation mechanisms through analyzing the type of negative ions on the oxidized surface. Analysis was performed using spectrometry mode with Bi3+ ions used for analysis and Cs (1 keV) ions used for sputtering on an analysis area of 350 μm2 and sputter area of 150 μm2.
2. Materials and methods 2.1. Preparation of substrate and coatings In this work, titanium aluminide was used as the substrate. The chemical composition of the substrate is presented in Table 1 and comprises of Ti and Al with small additions of Cr and Nb. Fig. 1 shows the XRD pattern of the titanium aluminide substrate. The substrate is composed of a mixture of TiAl and Ti3Al phases. Surface pretreatment of substrate was performed by grinding the substrate with SiC papers and cleaning with acetone to remove the surface oxides and/or contaminations before laser cladding. NiCr–Cr3C2 powder mixture was synthesized by mechanical alloying (MA) and rounded by spray drying in order to obtain spherical particles due to better flow properties [15]. The coaxial laser cladding process was performed in an argon environment using a powder feed system with a 0.7 KW transverse-flow fiber laser system equipped with a 4-axes CNC machine. Argon has been used as the shielding and powder carrier gas with the flow rate of 150 mL/min. For all experiments, the laser-beam spot size has been fixed of 2 mm. A constant laser power of 400 W with different scanning speeds (2–5 mm/s) and powder feeding rates (200–400 mg/s) were used to clad the as-milled NiCr–Cr3C2 powder mixture on the TiAl substrate to form single tracks. The final overlay for the proposed coating was achieved using a side-by-side overlapping (50%) of single track cladding with the optimal condition.
3. Results and discussion 3.1. Synthesis and characterization of NiCr–Cr3C2 powder mixture by mechanical alloying As mentioned before, Ni and Cr elemental powders were milled up to 40 h to prepare Ni–Cr solid solution. Fig. 2a shows the XRD patterns of Ni-20 wt% Cr powder mixture in as-mixed condition and after different milling times up to 40 h. The XRD pattern of as-mixed powder mixture consisted of Ni and Cr elemental peaks. Enhancement of the milling time to 10 h resulted in the reduction of peaks intensity and an increase in the peak width. This could be related to the crystallite size reduction and internal stress increment of the as-received elements [16–18]. Furthermore, Ni peaks shift slightly to the lower diffraction angles with increasing the milling time from 5 h to 30 h, confirming the dissolution of Cr atoms in the Ni crystal structure and formation of Ni (Cr) solid solution over time. The position of Ni peaks remain constant with further milling after 30 h indicating that further milling could only reduce the crystallite size of Ni(Cr) solid solution. The crystallite size of the nanostructured Ni(Cr) after 30h of milling was predicted as 43 nm using Williamson-Hall equation [19]. TEM analysis was also performed
Table 1 Chemical composition (wt.%) of the titanium aluminide substrate. Element
Ti
Al
Cr
Nb
Others
Chemical composition (wt.%)
60.54
32.51
2.27
4.65
0.03
2
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Fig. 1. XRD pattern of titanium aluminide substrate.
with Cr3C2. Chromium and graphite powders were also mixed with the molar ratio of 3:2 and milled up to 40 h in order to synthesize the Cr3C2 compound. Fig. 3 shows the XRD patterns of as-received and milled Cr–C powder mixture. Reduction of the intensity of chromium and graphite peaks are observed during mechanical alloying for up to 7 h.
to observe the microstructure of the formed Ni(Cr) compound after 30 h milling time (Fig. 2c). Nanometric grains are clearly visible in the individual powder particles shown in Fig. 2c. More information about the synthesis mechanism of the Ni(Cr) solid solution can be found elsewhere [20]. Since the dissolution of Cr in Ni is completed after 30 h of MA, the 30 h milled Ni(Cr) powder was used for subsequent mixture
Fig. 2. (a) XRD patterns of Ni-20 wt% Cr powder mixture at different milling times up to 40 h, (b) Magnified X-ray diffraction (XRD) peak around 2θ = 50–54, (c) TEM micrograph of the nanostructured Ni(Cr) solid solution after 30 h. 3
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Fig. 3. XRD patterns of Cr–C powder mixture with the molar ratio of 3:2 at different milling times.
powder mixture. The NiCr-50 wt.% Cr3C2 powder mixture was then rounded with spray drying to obtain spherical particles with the average size of around 50–120 μm, resulting in better flow during powder fed laser cladding.
By increasing the milling time to 10 h, a Cr3C2 peak appears beside the Cr peaks. Interestingly, further milling resulted to the formation of Cr7C3 instead of Cr3C2. Therefore, the XRD pattern of 40 h milled chromium-graphite powder mixture shows only Cr7C3 diffraction peaks. The same behavior (formation of Cr7C3 and Cr23C6 instead of Cr3C2) is also observed by Gomari et al. [21] during the milling of chromium-graphite powder mixture with the molar ratio of 3:2. Thermodynamic analysis was performed to understand the origin of such phenomenon. According to Equations (1) and (2) [21], the formation of both Cr3C2 and Cr7C3 chromium carbide compounds from the elemental powders is possible during mechanical alloying due to negative Gibbs free energy of formation. It is worth mentioning that the Gibbs free energy change of Cr7C3 at the milling temperature is more negative than that of Cr3C2. Graphite is a very ductile material with high tendency to stick to the milling media and balls during mechanical alloying. This leads to small changes of the molar ratio of Cr: C, resulting in the formation of Cr7C3 instead of the desirable Cr3C2 compound.
3Cr + 2C → Cr3 C2 ΔG (Cr3 C2) = −83189 − 0.967T J /mol ΔH298 = −72.3 kJ / mol
(1)
7Cr + 3C → Cr7 C3 ΔG (Cr7 C3) = −142666 − 30.92T J /mol ΔH 298 = −144.4 kJ / mol
(2)
3.2. Characterization of laser cladded coatings Fig. 5 shows the SEM micrographs of the single track cross-section of laser cladded coatings produced at laser power of 400 W with different scanning speed (S) and powder feeding rate (F). At a first glance, it can be seen that the geometrical properties of single track laser cladded coatings are highly dependent on the process parameters. No successful claddings were produced at higher powder feeding rate and low scanning speed due to detachment from the substrate (Fig. 5). Meanwhile, the powder-fed laser cladding process for all other combinations of scanning speed and powder feeding rate resulted in the formation of a sound cladding with unique geometrical properties. It should be mentioned that the overall heat input increases with the reduction of scanning speed, resulting in a high instability in the melt pool that is detrimental for the formation of high quality clad. Fig. 5 shows a decrease in clad height with increasing scanning speed and decreasing powder feeding rate. At the constant power of 400 W, the width of claddings was decreased with increasing the scanning speed. The reason can be explained by the reduction of the interaction time between the powder particles and laser beam at higher scanning speed [22]. In this case, the generated heat input to the centre of the molten pool decreases and as a consequence, the width of the claddings decreases, too. The enhancement of the powder feeding rate can also absorb some of the laser energy and results to lower heat input and therefore smaller clad width [22]. It can be observed that some of the coatings contain numerous porosities in their cross-section. Considering geometrical requirements and porosity content of the coatings, the sample cladded with the scanning speed of 3 mm/s and powder feeding rate of 300 mg/s was selected as the optimal cladding and considered for further characterization. It is worth to mention that the optimization of the laser cladding parameters for NiCr–Cr3C2 on TiAl substrate is comprehensively described in Ref. [22]. The surface of the optimal single track was studied to measure the surface roughness (Ra) (Fig. 6a-b). Roughness is an important cladding parameter which can influence the fatigue properties of the coating [23,24]. The fatigue mechanism can be divided into three stages: crack
As mentioned before, the milled Cr–C powder mixture is composed of the elemental peaks up to 7 h of milling time. Therefore, a two-step procedure consisting of 7 h milling and 3 h annealing at 900 °C was employed in order to fabricate Cr3C2. Fig. 4a shows the XRD patterns of 7 h milled Cr–C powder mixture after subsequent annealing. Cr3C2 diffraction peaks observed in Fig. 4a clearly show the presence of the more desirable Cr3C2 compound due to the reaction of activated chromium and graphite during annealing. Fig. 4b shows the TEM micrograph of the synthesized Cr3C2 after annealing. EELS analysis was employed in order to quantify the concentrations of Cr and C in the formed compound. It was found that the synthesized chromium carbide after annealing is composed of ∼40 at.% C which matches with the nominal composition of Cr3C2. The synthesized Ni(Cr) and Cr3C2 powder mixtures were mixed with equal weight ratio and further milled for 1 h to obtain a homogenous 4
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Fig. 4. (a) XRD patterns of Cr–C powder mixture after 7 h of milling and subsequent annealing at 900 °C for 3 h (b) TEM micrograph of the synthesized Cr3C2 after annealing.
initiation, cracks propagation and final failure. The first stage, crack initiation, plays an important role in low stress fatigue samples. It is well known that the rough surfaces normally exhibit poor fatigue properties due to the role of surface defects on the cracks initiation [23,24]. The average surface roughness of the single track surface was measured as 3 μm which seems to be sufficient for most practical applications. X-ray computed tomography (CT) was also performed on the coating and shows that the optimal cladding contains only 0.2% porosity with a median pore size of less than 20 μm (Fig. 6c). The microstructure of different parts of the single track cladding
was carefully studied by EBSD. Fig. 7 shows the SEM micrograph and EBSD result of the interface of titanium aluminide (TiAl) substrate and NiCr-chromium carbides single track coating. A thin layer of Ti3Al with the average thickness of ∼8 μm is formed between the substrate and cladding and acts as the interface. The microstructure of the coating close to the interface is very complex and is composed of different phases (Fig. 7b). Presence of different chromium carbide phases (Cr23C6 and Cr7C3) are observed close to the interface. In addition, the presence of TiC, Ti8C5 and Al4C3 close to the interface is probably caused by the dilution effect of the substrate. As the coating is applied by laser
Fig. 5. SEM micrographs of the cross section of the single-track laser cladded NiCr–Cr3C2 powder mixture with the laser power of 400 W and different scanning speed (S) and powder feeding rate (F). 5
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Fig. 6. (a) Surface profile of the optimal NiCr-chromium carbides single track cladding before and after curvature corrections (a–b) along with the result of computed tomography (CT) scan (c).
cladding, the surface layers of the substrate melt and elements from the substrate dilute in the cladded layer. The microstructure of the optimal cladding far away from the interface was also studied which shows a mixture of γ solid solution Cr7C3, Cr23C6 and the un-melted Cr3C2 particles having an average size of ∼10 μm (Fig. 8). Although, the as-prepared powder mixture has the chemical composition of NiCr-50 wt% Cr3C2, different carbide phases were detected on the cross-section of optimal cladding. One possibility for such behavior might arise from the partial melting of Cr3C2 during laser cladding and the formation of non-equilibrium chromium carbides upon solidification. The formation of non-equilibrium phases during the laser cladding of Cr3C2–NiCr feedstock powder have been also reported by Yang et al. [14]. Complete transformation of Cr3C2 to Cr7C3 was reported by Yang et al. [14]. Heubner et al. [25] has also reported partial dissolution of WC fine particles during laser cladding of IN625WC composite powder, resulted to the formation of different secondary carbides with higher amounts of Nb, Cr and W. It is worth mentioning that, in this investigation, some traces of un-melted Cr3C2 was observed mostly on the regions close to the top surface of the cladding which is the last region of solidification. In fact, the Cr3C2 particles are partially melted and solidified to Cr7C3 and Cr23C6 and the remaining un-melted particles floated on the melt and remain there. It should be noted that TiC, Ti8C5 and Al4C3 were not observed away from the substrate. Nanoindentation tests were performed to assess the effectiveness and performance of the cladded layer. Fig. 9 shows the load-displacement nanoindentation curve of the substrate and the optimal cladding. Since, the nanoindentation test was conducted in load control mode, lower penetration depth of the cladding reflects its higher hardness compared to the titanium aluminide substrate. The hardness and elastic
Fig. 7. (a) SEM micrograph of the interface of substrate and the optimal single track coating (b) EBSD phase map of the area close to the interface. 6
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H 2 ⎛1 E ln ⎜ tanα ⎞⎟ =K+ σy 3 ⎝ 3 σy ⎠
(3)
where H and E represent the hardness and elastic modulus obtained from nanoindentation test, σy is the yield strength, α is a constant angle defined as 19.7° for Berkovich indenter and K = 4/3. Other values of K by Studman and Bushby [28], were reported as 7/6 and 1.15. Fig. 9 shows the yield strength of the substrate and the optimal laser cladded coating based on different cavity models [27,28]. The value of yield strength is in the same range for all models, about 1.1 GPa for TiAl and 3.6 GPa for the optimal cladding. Based on the results from nanoindentation tests and yield strength, it can be concluded that the mechanical behavior of the substrate is improved by laser cladding. Since the final cladded surface layer forms from the overlapping of single tracks, its microstructure was carefully studied to correlate it to the mechanical properties oxidation resistance. Fig. 10a shows the cross-section of the overlapped NiCr-chromium carbides coating on titanium aluminide substrate. After overlapping, the average thickness of the coating increases to ∼500 μm. SEM analysis reveals presence of voids in the overlapped coating. A detailed CT analysis of the overlapped coating showed a 2.3% porosity and is presented in Fig. 10b. CT results show a higher porosity than the optimal cladded layer but a lower porosity than thermal spray deposition techniques such as Air plasma spray (APS) [29] and high-velocity oxy fuel (HVOF) [30,31]. Repeated melting of the coating during overlapping may cause turbulence in the melt pool, resulting in a higher amount of trapped gasses during solidification [32,33]. It was observed that most of the porosities have a radius less than 40 μm and only limited number of pores have larger radii. The XRD pattern of the overlapped coating is shown in Fig. 11 and confirms that the coating is composed of a mixture of γ and different chromium carbides (Cr3C2, Cr7C3 and Cr23C6). 3.3. Evaluation of oxidation resistance
Fig. 8. (a) SEM micrograph, (b) EBSD phase map of the NiCr-chromium carbides single-track cladding at the middle height of coating.
Fig. 12 shows the specific mass gain (mass gain per unit area) versus oxidation time for the substrate and overlapped coating at 900 °C. As seen, the specific mass gain of the coating is lower than the substrate. The specific mass gain of the laser cladded coating is ∼0.004 mg/cm2, twice lower than ∼0.009 mg/cm2 for the substrate. It should be noted that, for both samples, the slope of the mass gain curve is always higher at the beginning of oxidation and reduces with increasing oxidation time. To quantify the rate of oxidation, the curve of (mass gain per unit area)2 versus oxidation time was plotted for both substrate and coating and is presented in Fig. 12b. As can be observed, a straight line fits the data points from TGA experiments, meaning that the oxidation kinetics is parabolic. The value of R2, for the experimental data of substrate and
modulus of NiCr-chromium carbides cladding were recorded as 9.5 ± 1.2 GPa and 250 ± 11 GPa, respectively which is higher than that of TiAl substrate (hardness = 3.6 ± 0.2 GPa and elastic modulus = 164 ± 4 GPa). The presence of different types of chromium carbides in the microstructure of the optimal cladding is responsible to the remarkable increase of the hardness of the coating, due to the higher hardness of carbide particles compared to the Ni and/or NiCr alloys [26]. Johnson cavity model [27] was used to estimate the yield strength of the substrate and cladding from nanoindentation data (Equation (3)).
Fig. 9. Typical load-displacement curve of TiAl substrate and NiCr-Chromium carbides optimal cladding during nanoindentation test under load-control mode and yield strength (GPa) of the TiAl substrate and NiCr-chromium carbide single track cladding estimated by cavity models. 7
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Fig. 10. (a) Cross section overview of NiCr-chromium carbides overlapped coating produced by laser cladding and (b) its related computed tomography (CT) scan showing the porosity size.
compared to the substrate. It is reported that Cr2O3 is stable up to the temperature of 950 °C, and therefore can protect the surface coating at high temperatures [37]. There are some minor peaks related to different carbides and also γ solid solution which might be a result of the original coating. It is reported that the thickness of oxide film on the surface of Cr3C2-25 wt%NiCr is less than 2 μm after 16 days of oxidation at 800 °C and do not form a continuous surface layer before 8 h of oxidation [38]. Therefore, the peaks related to γ and different chromium carbide phases from the XRD pattern of oxide layer might be generated due to the effect of original coating. The formation of Cr23C6 carbide phase is also previously reported [12] during isothermal oxidation of Cr3C2-25% NiCr at 800 °C and attributed to the stepwise oxidation of Cr3C2. The SEM morphology of the coating surface after oxidation test also confirms that the formed layer is not continuous and some parts of the original coating are still visible on the surface. SIMS studies were carried out to study the oxidation mechanisms of the substrate and coating. Fig. 15a–b shows the SIMS depth profiles of titanium aluminide and NiCr-chromium carbides oxidized at 900 °C for 5 h, respectively. Results from the titanium aluminide substrate show that the intensity of TiO− and TiO2− is high at the surface of oxide, while decreasing continuously across the depth and instead the intensity of AlO− and AlO2− negative ions is grown. It further confirms that the oxidation of titanium aluminide starts with the formation of Al2O3 and then continues with the formation of TiO2 which is in good agreement with the results of the XRD and EDS measurements on the top surface of the oxide layer. It is worth to mention that the crystal lattice of TiO2 contains both oxygen vacancies and interstitial titanium atoms. The former prevails at low temperatures and high oxygen pressure and the latter at high temperatures and low oxygen pressure. It is reported that at T < 800 °C, TiAl is mainly oxidized through diffusion of oxygen from the oxidizing media towards the interface of substrate/oxide layer [35]. However, the oxygen partial pressure of inner
coating was calculated to be 0.94 and 0.985, respectively. A parabolic oxidation rate can be represented by Equation (4), where “t” is oxidation time and K and C are the kinetics constants [34]. (Δm/a)2 = kpt + C
(4)
The slope of fitted straight lines (parabolic rate constant-Kp) for substrate and coating is then calculated and given in Fig. 12b. The parabolic rate constant of the coating is 5 times smaller than the substrate indicating the improved oxidation resistance of substrate after laser cladding. The morphology of oxide layer formed on the surface of titanium aluminide after 5 h of oxidation is shown in Fig. 13(a–b). A dense layer of oxidation products was formed on the surface of substrate after oxidation. The GAXRD pattern reveals that the outer surface of the oxide layer is mostly composed of TiO2 with some minor amount of Al2O3 (Fig. 13c). EDS analysis on the surface of oxide layer also shows that the oxide layer is composed of ∼60 at.% O, 36.5 at.% Ti and 2.5 at. % Al with small additions of Nb and Cr. These results confirm that the outer surface of the oxide layer is almost fully covered with TiO2 (Rutile). It is reported that the affinity of Ti and Al to oxygen is very close [35]. In the other words, both TiO2 and Al2O3 have similar thermodynamic stability. However, the growth rate of TiO2 is faster than Al2O3 due to its unique crystal structure which provides extra sites for diffusion of oxygen. However, the oxide layer is composed of both Al2O3 and TiO2 at the early stages of oxidation, TiO2 nuclei grow faster and the outermost layer of oxide layer is enriched with rutile [36,37]. The morphology and XRD pattern of oxide layer formed on the surface of coating after oxidation test is shown in Fig. 14. Results in Fig. 14 show that chromium oxide (Cr2O3) is the major phase in the oxide layer. The formation of chromium oxide-rich scale in the present study probably acts as a barrier to the diffusion of oxygen into the coating and therefore improved oxidation resistance of coating
Fig. 11. XRD pattern of NiCr-chromium carbides overlapped coating. 8
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Fig. 12. (a) Mass gain per unit area and (b) (mass gain per unit area)2 versus oxidation time for TiAl substrate and NiCr-chromium carbides coating during TG experiments at 900 °C and parabolic rate constant of substrate and coating in high temperature oxidation test at 900 °C.
2 The microstructure of the coating was mainly composed of γ and different chromium carbide phases (Cr3C2, Cr7C3 and Cr23C6). The formation of different chromium carbide phases might arise from the partial melting of Cr3C2 powders and formation of non-equilibrium Cr7C3 and Cr23C6 phases. 3 The NiCr-chromium carbides coating showed a superior oxidation resistance at 900 °C compared to the titanium aluminide substrate. 4 The oxidation behavior of NiCr-chromium carbides coating and titanium aluminide substrate follow the parabolic law. The parabolic rate constant of the coating was one fifth of the substrate at 900 °C. 5 ToF-SIMS analysis revealed that the oxidation of substrate started with the formation of Al2O3 and further continued with growth of TiO2. In contrast, the surface of oxide layer on the NiCr-chromium carbide coating was almost fully covered with Cr2O3.
layers decreases with increasing the temperature and/or thickness of the oxide layer, resulting to the high concentration of titanium interstitials which promotes further diffusion of Ti towards outer layers [35]. In the case of NiCr-chromium carbide coating (Fig. 15b), it can be observed that the intensity of CrO−, CrO2− and CrO3− is very high at the surface and seems to be constant throughout the thickness of the oxide layer. It is reported in literature that the CrO− ions provide the best clue for the presence of Cr2O3 in the oxidized region [39]. Therefore, it can be concluded that Cr2O3 is the main compound in the NiCr-chromium carbide coating. 4. Conclusions In the present study, the NiCr–Cr3C2 powder mixture was synthesized by mechanical alloying and then coated on the TiAl substrate using the powder-fed laser cladding process. The following conclusions can be drawn from this study:
Acknowledgements The authors would like to thank Canadian Centre for Electron Microscopy (CCEM) and Department of Mechanical Engineering at McMaster University for conducting the EBSD, TEM and Nanoindentation measurements. The authors gratefully thank WATLAB at the University of Waterloo for conducting the SIMS experiments. The authors would also like to thank the members of Multi-Scale Additive
1 The optimized cladding was obtained at P = 400 W, S = 3 mm/s and F = 300 mg/s that resulted in single tracks with approximately 0.2% porosity. The final overlay for oxidation test was achieved through the side-by side overlapping (50%) of the single track cladding with the optimal laser condition. 9
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Fig. 13. (a–b) Morphology and (c) GAXRD pattern of oxide layer formed on the surface of titanium aluminide after high temperature oxidation test at 900 °C for 5 h.
Fig. 14. Morphology and GAXRD pattern of oxide layer formed on the surface of NiCr-chromium carbides coating after high temperature oxidation test at 900 °C for 5 h. 10
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Fig. 15. SIMS profile of the (a) titanium aluminide substrate and (b) NiCr-chromium carbide coating after high temperature oxidation test.
Manufacturing (MSAM) group especially Prof. Mihaela Vlasea, Dr. Vladimir Paserin, Dr. Hamed Asgari and Dr. Yahya Mahmoodkhani for their supports.
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