Developing the ductility and thermal fatigue cracking property of laser-deposited Stellite 6 coatings by adding titanium and nickel

Materials and Design 162 (2019) 271–284

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Developing the ductility and thermal fatigue cracking property of laser-deposited Stellite 6 coatings by adding titanium and nickel Ying Wu, Yan Liu ⁎, Hui Chen ⁎, Yong Chen, Hongyu Li, Xinyu Cao School of Materials Science and Engineering, Southwest Jiaotong University, Key Laboratory of Advanced Technologies of Materials, Ministry of Education of China, Chengdu 610031, PR China

H I G H L I G H T S

G R A P H I C A L

A B S T R A C T

• A strategy is given to improve the ductility and thermal fatigue cracking feature of Stellite 6 coating by adding Ti and Ni. • Ti addition made the M23C6 eutectic carbide change to the TiC carbide; Ni addition stabilized the γ-Co in matrix. • 1.0Ti20Ni coating exhibited the highest elongation to fracture and the lowest thermal fatigue crack propagation rate.

a r t i c l e

i n f o

Article history: Received 13 September 2018 Received in revised form 29 November 2018 Accepted 30 November 2018 Available online 4 December 2018 Keywords: Stellite 6 coating Coating design Microstructure Mechanical properties Thermal fatigue crack propagation

a b s t r a c t A design strategy is presented for the development of a laser-deposited Stellite 6 coating with high ductility and low thermal fatigue crack propagation rate. The strategy involved adding titanium to modify carbide morphology, accompanied by an increase in the nickel content to stabilize the γ-Co phase in the matrix. The molar ratios of the titanium to the carbon in coatings were 0.5, 1.0, 1.5, and 2.0. The amount of nickel added to the titaniummodified coatings was 10 and 20 wt%. The net-like M23C6 eutectic carbides in the Stellite 6 coating transformed to isolated TiC particles with increasing molar ratio of titanium to carbon. Adding nickel to the coatings resulted in an increase in the stacking fault energy and the stabilization of the γ-Co phase. Accordingly, the ductility increased with increasing nickel concentration. The modified coating, named 1.0Ti20Ni, in which the molar ratio of titanium to carbon was 1.0 and the nickel concentration was 20 wt%, exhibited the highest elongation to fracture (~12%), along with a high strength (1230 MPa) in all coatings, Besides, the 1.0Ti20Ni coating revealed the lowest thermal fatigue crack propagation rate. © 2018 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http:// creativecommons.org/licenses/by-nc-nd/4.0/).

1. Introduction Thermal fatigue cracking is one of the main life-limiting factors for the brake discs of trains. Brake discs [1–3], which play an important role in speed reduction and regulation through frictional action [4], transform the kinetic energy of vehicles into thermal energy. This causes ⁎ Corresponding authors. E-mail addresses: [email protected] (Y. Liu), [email protected] (H. Chen).

a great temperature fluctuation on the worn surface of a brake disc. In the case of emergency braking of a train travelling at the speed of 300 km/h, the highest temperature observed on the worn surface can exceed 600 °C [5]. The huge temperature gradient and fluctuation may produce a local stress concentration, leading to the initialization of cracks. These cracks will propagate during subsequent braking, resulting in premature failure of the brake disc. Therefore, the thermal fatigue crack propagation rate is the criterion used to evaluation the potential of materials for brake discs. Currently, Cr-Mo-V and Cr-Ni-Mo-V

https://doi.org/10.1016/j.matdes.2018.11.063 0264-1275/© 2018 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).

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low-alloy casting steels are mainly employed to fabricate the brake discs used in trains owing to their favorable strength and toughness [6]. As structural materials of brake discs that transfer friction force to the axle (note that brake discs are mounted on an axle through flanges), Cr-Mo-V steel and Cr-Ni-Mo-V steel are excellent. However, their poor elevated temperature strength and oxidation resistance accelerate the nucleation and propagation of cracks during the thermal fatigue process [7]. Consequently, both Cr-Mo-V steel and Cr-Ni-Mo-V steel are not competent enough to be used as materials for the worn surface layer of brake discs. Based on the service condition, it is desirable that the various parts of brake discs display different properties, e.g., the substrate has high strength and toughness, whereas the worn surface exhibits good thermal fatigue performance. Hardfacing is an optional technique used to improve the wear resistance and elevated temperature performance of components exposed to severe service conditions. It is also a flexible technique that can be accomplished through a variety of methods, such as laser deposition, plasma transfer arc welding and tungsten inert gas arc welding [8–10]. Compared with the other methods, laser deposition is favored because of its high efficiency, limited heat input, small distortion, low dilution rate, and the high bonding strength with the substrate. In addition, the feed materials for laser deposition can be indifferent forms, including powder, wire, and strip. In particular, for powder feed materials, it is convenient to adjust the composition of the powder through mechanical mixing and atomization. Therefore, laser deposition is a good choice for depositing hardfacing materials on brake discs. Based on the braking condition of the brake discs used in trains, cobalt-based alloys are potential hardfacing materials that are used as the worn surface layer owing to their excellent wear resistance, heat resistance, thermal fatigue property, and corrosion resistance at elevated temperatures [11–13]. Stellite alloys are the typical Co-Cr-W alloys. Among them, Stellite 6 is widely used in industrial applications as a hard surface layer. Its high carbon and chromium contents promote carbide precipitation, which provides it with high hardness and wear resistance. Chromium also improves hot corrosion resistance [14,15], whilst other refractory metals (e.g., molybdenum and tungsten) are added for solid solution strengthening [16]. A certain content of nickel stabilizes the γ-Co matrix that is face-centered cubic structure. Besides, Stellite 6 exhibits favorable thermal fatigue and creep properties [16]. Our research group prepared a Stellite 6 coating on a brake disc by using laser deposition and evaluated its thermal crack propagation rate. However, the obtained results could not satisfy the requirement that the crack propagation rate of the Stellite 6 coating should be lower than that of the low-alloy casting steel currently used in brake discs. Although most of the references report positively on the properties of Stellite 6 coatings, it seems that the ability of the coating to resist thermal crack propagation is not sufficient, which was also observed by Birol [17,18]. In further studies, the mechanism of thermal cracking of Stellite 6 coating was revealed. The oxidized net-like M23C6 (eutectic carbides) and the interface of γ-Co and ε-Co acted as pathways to promote crack propagation during thermal fatigue. These results were similar to that obtained by Tang and Tunthawiroon [19,20]. Based on the failure mechanism disclosed above, it is believed that modifying the morphology of the carbides, which should be changed from the netlike shape to isolated particles, and stabilizing the γ-Co phase are the key to reducing the thermal fatigue crack propagation rate. In order to modify the eutectic carbides, the precipitation temperature of the carbides should be increase and the eutectic reaction should be suppressed. Titanium is a strong carbide-forming element [21]. Titanium carbide (TiC) has high melting temperature (3147 °C), thermodynamic stability, and wear resistance [21–23]. Shahroozi et al. [24] deposited Stellite 6/ TiC mixed powder on a steel substrate. They found that the eutectic structure was eliminated when the content of TiC was 10 wt%. Ma et al. [25] added titanium to the Ni60/WC coating and found that the TiC precipitated around Cr5B3 instead of the M23C6 (eutectic carbides). From the results listed above, it can be noted that whether the additive

was TiC or pure titanium powder, the types and morphologies of carbides changed in coatings. The chromium in Stellite 6 is another strong carbide-forming element and is a major constituent element of M23C6. Ma et al. [25] pointed out, theoretically, M23C6 was more prone to form in the equilibrium state than TiC. However, in the study of Huang et al. [26], in-situ TiC particles reinforced nickel-based composite coatings were produced by plasma spray welding with mixed powders (NiCrBSi + Ti + NiCr–Cr3C2). The results showed that the TiC in situ precipitated first from the molten pool earlier than the chromium carbides. Therefore, pure titanium powder and TiC particles could suppress the formation of M23C6 in hardfacing process. Considering the service conditions of brake discs, high hardness is not necessary. The addition of TiC increases the hardness of a coating, but easily induces cracking [24]. Pure titanium powder is the first choice. To stabilize of the γ-Co phase, the γ/ε martensitic transformation should be suppressed. This transformation results from a change in the stacking sequence of the close-packed {111}γ planes from the ABCABC to the ABABAB structure [27], which is associated with the stacking faults energy (SFE) [28]. Nickel is generally added to stabilize γ-Co because it could increase the SFE [28,29]. Li [30] pointed out that the amount of nickel did not influence the presence of the TiC synthesized in situ. Seldom research gave the reposts about that titanium could promote the ε-Co formation [31]. Therefore, the additions of titanium and nickel may be the potential choice to modify Stellite 6 coatings. The modified Stellite 6 coating with improved performance may play an important role in prolonging the service life of brake discs. However, limited reports can be found to study the thermal fatigue cracking property of a Stellite 6 coating modified with both titanium and nickel [31–33]. In the present study, a design approach was presented to obtain cobalt-based coatings with good ductility and low thermal fatigue crack propagation rate that were fabricated on brake discs by laser deposition. For this purpose, the simultaneous addition of pure titanium and nickel powders to the original Stellite 6 powder was employed. The mixed powders, containing different amounts of titanium and nickel were deposited on Cr-Ni-Mo steel. The dependence of the microstructure and mechanical properties on the titanium and nickel concentrations was examined to optimize the chemical compositions. 2. Materials and methods 2.1. Coating design The basic design strategy for the titanium- and nickel- modified Stellite 6 coatings used in brake disc applications is briefly described, along with the results of thermodynamic calculations. The phase diagrams were constructed by using the Pandat software with ferrous database. Fig. 1 shows the vertical sections of the phase diagrams calculated for the Co-27Cr-1C-Ti, Co-24Cr-11Ni-0.9C-Ti and Co-24Cr21Ni-0.9C-Ti systems. The proportions of components in the phase diagrams represent the chemical compositions of the coatings that were modified in this study. As seen in Fig. 1a, with the addition of titanium, the MC carbide precipitated for the titanium concentration range ~0.6 to 10.0 wt% after the liquid phase solidification, which indicated that titanium could modify the type of carbides in original Stellite 6. M7C3 also precipitated for titanium concentrations less than ~4.7 wt%, and gradually transformed to M23C6. The eutectic structure consists of γ-Co and carbides that were formed at ~3.1 wt% titanium when the temperature was about 1233 °C. The amount of titanium added was the key aspect affecting the types of carbides in the coating. Nickel was selected as an additive to improve the ductility, as increasing the nickel concentration could stabilize γ-Co. By comparing the phase diagrams of Fig. 1a, b, and c, it is noted that the addition of nickel lowered the formation temperature of the ε-Co phase (marked with red lines) and enlarged the phase regions containing the γ-Co phase. The addition of nickel also effectively stabilized the γ-Co phase by increasing the SFE, which suppressed the athermal γ-Co to ε-Co martensitic transformation that occurred during

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Fig. 1. Vertical section of calculated phase diagrams of (a) Co-27Cr-1C-Ti, (b) Co-27Cr-11Ni-0.9C-Ti, and (c) Co-27Cr-21Ni-0.8C-Ti systems obtained using the Pandat software.

cooling [28,34,35]. In addition, nickel did not alter the constituent phase when the temperature was higher than 900 °C. Here, it should be noted that laser deposition is a process in which an extremely high cooling speed, 104–106 K/s, can be obtained [36], along with nonequilibrium solidification, which was accompanied by serious component segregation [37,38]. The phases of the laser-deposited coatings were usually in the nonequilibrium state. The phase diagrams that were constructed based on equilibrium thermodynamic calculations could not precisely predict the constituent phases of the coatings at room temperature. However, the trend in the phase transformation observed in the phase diagrams might shed light on the design of the titanium- and nickelmodified Stellite 6 coatings. From the discussions above, it is noted that the addition of titanium and nickel had positive effects on the modification of the carbide morphology and the stabilization of the γ-Co phase. 2.2. Sample preparation The powder mixtures used in this study comprised commercial Stellite 6 powder (supplied by Höganäs, Sweden) and pure titanium and nickel powders (both 99.9 wt% purity, supplied by Beijing AMC Powders, China). All of them were spherically atomized powders having diameters between 80 and 150 μm. The pure titanium and Stellite 6 powders were blended such that the molar ratios of titanium to the carbon in mixed powders were 0.5, 1.0, 1.5 and 2.0. The coatings deposited by using these mixed powders are hereinafter referred to as 0.5Ti, 1.0Ti, 1.5Ti and 2.0Ti, respectively. Based on the microstructures of the titanium-modified coatings, the pure nickel powder was further added to levels of 10 and 20 wt%, except the case of the 0.5Ti powder because of the existence of a eutectic structure (discussed in Section 3.1). The corresponding coatings with nickel added are referred to as 1.0Ti10Ni, 1.0Ti20Ni, 1.5Ti10Ni, 1.5Ti20Ni, 2.0Ti10Ni and 2.0Ti20Ni. All the powders were mixed mechanically by using a planetary ball milling machine (YXQM-2L) for 3 h at the rotating rate of 200 rpm. ZrO balls with diameters ranged from 4 to 10 mm were used. The volume ratio of ball to powder was 1.0. The equivalent chemical composition (in weight ratio of the mixed powders) is listed in Table 1. The laser deposition process was performed on an IPG YLS-4000 laser system. The system was equipped with a fiber-delivered 4.0 kW solid state laser and a coaxial laser deposition head with a focal length of 300 mm. Argon was used as the powder carrier and shielding gas. 24CrNiMo steel, used in brake discs, was employed as the substrate, and its chemical composition is listed in Table 1. The optimum processing parameters are summarized in Table 2. No cracks or porosity was observed. In order to satisfy the dimensional requirement of the specimens, the coatings were fabricated in three layers. The orthogonal deposition direction was adopted to minimize the internal stress, as shown in Fig. 2a.

2.3. Microstructural characterization The microstructures were examined by using a field-emission scanning electron microscope (FE-SEM; FEI, XL30S-FEG) at the acceleration voltage of 20 kV. Energy dispersive spectroscopy (EDS; Oxford, X-MaxN) was used to determine the chemical compositions of the phases. X-ray diffraction (XRD) was carried out by using a Pert-Pro MPD diffractometer. Electron back scatter diffraction (EBSD. Oxford, NodlysNano) was performed with a FE-SEM (JEOL, JSM-7800F) by employing a 50 nm scanning step. Transmission electron microscopy (TEM) was performed by using a FEI TECNAI F20 instrument at 200 kV. Before XRD, SEM and EBSD examinations, the specimens were mechanically grinded and mirror polished. The samples for SEM were etched by electrochemical corrosion in a solution of 1/9 (vol%) hydrochloric acid/ alcohol at 3.5 V. The EBSD samples were prepared by ion polishing with Gatan llion 697 to obtain the required surface quality. The samples for TEM were sectioned to 0.2 mm thickness by electric discharge machining, thinned by grinding, and then ion milled by Gatan PIPS 695.

2.4. Mechanical characterization Uniaxial tensile tests were performed at room temperature on the specimens having the geometry shown in Fig. 2b. The tensile specimens were cut from the top two layers of the coatings by electric discharge machining. The loading direction was parallel to the direction of deposition of layer 3 and perpendicular to that of layer 2 (Fig. 2a). The tensile speed was 0.5 mm/min. An extensometer was used to measure the deformation in the gauge length. The fractured surfaces of the tensile test specimens were observed under a FE-SEM, and the cross-section microstructures near the fracture surfaces were characterized by EBSD.

Table 1 Chemical composition (wt%) of Stellite 6 powder and modified powder.

Stellite 6 0.5Ti 1.0Ti 1.5Ti 2.0Ti 1.0Ti10Ni 1.0Ti20Ni 1.5Ti10Ni 1.5Ti20Ni 2.0Ti10Ni 2.0Ti10Ni 24CrNiMo

C

Ti

Ni

Si

Fe

Cr

Co

W

Mo

Mn

1.10 1.08 1.05 1.03 1.01 0.95 0.84 0.93 0.82 0.91 0.81 0.24

/ 2.14 4.20 6.17 8.06 3.78 3.36 5.55 4.94 7.25 6.45 /

1.50 1.47 1.44 1.41 1.38 11.30 21.15 11.27 21.13 11.24 21.10 1.02

1.00 0.98 0.96 0.94 0.92 0.86 0.77 0.85 0.75 0.83 0.74 0.35

1.50 1.47 1.44 1.41 1.38 1.30 1.15 1.27 1.13 1.24 1.10 Bal.

28.50 27.89 27.30 26.74 26.20 24.57 21.84 24.07 21.39 23.58 20.96 0.73

Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. /

4.40 4.31 4.22 4.13 4.05 3.80 3.38 3.72 3.30 3.65 3.24 /

/ / / / / / / / / / / 0.54

/ / / / / / / / / / / 0.94

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Table 2 Parameters used for laser depositing. Laser power (W)

Laser spot size (mm)

Transverse speed (mm/min)

Powder flow rate (g/min)

Step-over width (mm)

Carrier gas flow (L/min)

Shielding gas flow (L/min)

1800

4.0

280

11.3

2.0

3

35

Thermal fatigue crack propagation tests were carried out by using plate specimens having the dimensions shown in Fig. 2c. These specimens were also cut from the top two layers of the coatings by electric discharge machining; the cutting was followed by grinding and mirror polishing. A notch was machined on the shot edge of the specimens to introduce a stress concentration, so that a crack can be initiated there. The thermal fatigue crack propagation tests were carried out in the form of cyclic heating and quenching, as the schematic diagram shown in Fig. 2d. In each thermal cycle, the specimens were first heated to 650 °C in a tube furnace for 200 s and then drowned in flowing water for 15 s to cool to room temperature. The highest and lowest temperatures observed in the test corresponded to the highest and lowest temperatures that the brake discs potentially encountered during emergency braking [2,7]. The parameter monitored during the thermal cycling was the temperature of the sample, which was measured by Ktype thermocouples that were fixed 2 mm away from the notch, as shown in Fig. 2d. The plot of temperature versus time is also shown in Fig. 2d. The crack length was measured by optical microscopy (OM; ZEISS, Observer A1m) after a certain number of thermal cycles. 2.5. Computational thermodynamics of the SFE The computation method of the thermodynamic SFE was based on Olson and Cohen [28,39,40] formulation: γ sf ¼ 2ρΔGγ→ε þ 2σ γ=ε

ð1Þ

where γsf is the SFE (mJ/m2), ΔGγ→ε is the Gibbs free energy change for the γ → ε phase transformation, and σγ/ε is the interfacial energy per unit area of the phase boundary. Based on the research of Tria [28,41], the value of σγ/ε was assumed to be 7.5 mJ/m2. A detailed description of the calculation process can be found in the report of Tria [41], and the modifications made in this study are listed in Appendix A. The thermodynamic data used in this study was the Scientific Group Thermodata Europe data for pure elements [42] and those found in published references. In this study, the cobalt matrix was considered as a Co-Cr-Ni-Ti quaternary alloy according to the major elements present. The effect of carbon segregation on the SFE, which was similar to the effect of nitrogen in austenitic steels [40], was not considered. This was because the carbon in Stellite 6 and the modified coatings mainly formed carbides, and underwent only limited dissolution in the cobalt matrix. In addition, the carbon concentration cannot be accurately

detected easily. Here, as Tria did [41], the ternary excess Gibbs free energy, EGi, j, kγ→ε, was introduced. However, limited data was available, only a partial ternary excess Gibbs free energy could be used, and no ternary magnetic parameter could be applied. 3. Results 3.1. Microstructure Fig. 3 shows the XRD patterns of the Stellite 6 and modified cobalt coatings. All the coatings primarily exhibited the peaks of γ-Co. The addition of titanium (for 0.5Ti, 1.0Ti, 1.5Ti and 2.0Ti) caused the weakness of peaks intensity of γ-Co, compared with Stellite 6, whereas the addition of nickel (for 1.0Ti10Ni, 1.0Ti20Ni, 1.5Ti10Ni, 1.5Ti20Ni, 2.0Ti10Ni and 2.0Ti20Ni) made the peaks intensity of γ-Co increased. The εpeaks were only identified in the titanium-modified cobalt coatings; they were not detected after nickel was added to the coatings. The addition of Ti resulted in the peaks of TiC at 2θ ~36° and ~42°, which are magnified on the right side of Fig. 3. Both M7C3 and M23C6, could be simultaneously identified in all the coatings, but their peaks overlap with the peaks of the γ-Co and ε-Co phases. It was therefore hard to determine exactly whether these carbides existed. Fig. 4 shows the SEM images of the Stellite 6 coating and the titanium-modified coatings with different molar ratios of titanium to carbon. As seen in Fig. 4a, fine dendrites and net-like interdendritic eutectics are observed in the basic microstructure of the Stellite 6 coating. The dendritic cobalt matrix was formed at the first stage of solidification. Then, a eutectic structure containing a cobalt-rich phase and the eutectic carbides solidified in the interdendritic region. A gradual change in the microstructure is observed upon increasing the titanium content (i.e., increasing the molar ratio of titanium to carbon). The eutectic structure (marked with yellow arrows in Fig. 4b) still remained in the 0.5Ti coating, but vanished in the 1.0Ti and 1.5Ti coatings (Figs. 4c and d). Isolated particles replaced the original net-like eutectic carbides. With further increase in the titanium content, some irregular bulk phases (marked with red arrows in Fig. 4e) appeared in the 2.0Ti coating. In addition, the 2.0Ti coating could easily crack during the laser deposition process, therefore, no tensile specimens or thermal fatigue cracking specimens could be prepared with this coating. The particles that precipitated in the titanium-modified coatings exhibited several shapes including flower-like, strip, polygon and irregular bulk, as shown in Figs. 4b to 4e. The EDS results of the eutectic carbide and

Fig. 2. Schematic diagram of specimens and thermal fatigue process: (a) orthogonal depositing direction of three layers coating, (b) specimen size of tensile test, (c) specimen size of thermal fatigue crack propagation test, and (d) thermal fatigue process.

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Table 3 Spectrum of EDS test points in Fig. 4 (mol %).

Sp1 Sp2 Sp3 Sp4 Sp5

Fig. 3. XRD patterns for Stellite 6 and modified cobalt coatings.

the differently shaped particles of the titanium-modified coatings (marked with numbered points) are listed in Table 3. The flower-like, strip, and polygon (Sp2, Sp3, and Sp4 in Fig. 4b) particles contained both titanium and chromium, compared to the eutectic carbide (Sp1 in Fig. 4a) rich in chromium. It is inferred that the particles with the flower-like, strip and polygon shapes were TiC-type carbide containing a certain amount of chromium (hereafter, this carbide is referred to as TiC). Furthermore, the flower-like particle contained oxygen. Besides carbon, titanium also has an affinity for oxygen, acting as a deoxidizing agent in welding materials to remove air holes. The irregular bulk particles (Sp5 in Fig. 4e), which were rich in cobalt, chromium and titanium, but contained less of carbon, were determined to comprise intermetallic

C

O

Si

Ti

Cr

Mn

Fe

Co

Ni

Mo

W

40.32 36.79 32.06 55.11 13.65

/ 7.36 / / /

1.44 / 1.09 0.33 3.92

/ 23.41 5.14 12.53 9.87

25.98 5.83 21.93 11.95 21.56

0.17 0.1 0.11 0.04 0.13

2.21 0.9 2.53 2.03 2.27

28.18 7.96 35.39 15.04 47.15

0.32 0.15 0.74 0.24 0.76

0.19 0.11 0.17 0.26 0.22

1.19 1.56 0.85 2.47 0.48

compounds. The exact types of these particles above were further characterized through EBSD and TEM, and the results are presented in the following sections. To calculate the fraction of precipitates (eutectic structure, TiC and intermetallic compound) of coatings (Stellite 6, 0.5Ti, 1.0Ti, 1.5Ti, and 2.0Ti), five randomly-selected microstructure SEM images for each coating was measured by Image-pro software to calculate the area fraction of precipitates, and the result was shown in Fig. 4f. The area fraction of the whole precipitate first decreased with increasing titanium content up to the 1.0Ti coating (to 4.2 wt% titanium concentration), and then increased to ~24 wt% for the 2.0Ti coating. The area fraction of TiC also increased with increasing Ti concentration up to the 1.5Ti coating (to 6.17 wt% titanium concentration), which corresponded to the gradual replacement of the eutectic structure gradually with TiC, whereas it decreased in the 2.0Ti coating as the intermetallic compound appeared. As the 0.5Ti coating still contained the eutectic structure, no nickel could be further added to it. On the other hand, the 1.0Ti, 1.5Ti and 2.0Ti coatings were added with 10 and 20 wt% of nickel. Fig. 5a to c and e to g show the SEM images of the coatings modified with both titanium and nickel. With the addition of nickel, the morphology of the particles become flower-like, strip, and polygon, similar to those of the particles in the titanium-modified coatings (Fig. 4a–e). No eutectic structure or intermetallic compounds were found. Five randomlyselected microstructure SEM images for each coating was measured by Image-pro software to calculate the area fraction of precipitates, and the result was shown in 5d. The area fraction of the precipitates decreased with increasing nickel concentration. The average diameter of the TiC particles in each of the modified coatings, measured by Imagepro software, was similar, as revealed in Fig. 5h. It seems that the geometry of TiC was not affected by the concentration of titanium and nickel employed in this study. Fig. 6 shows the SEM images and the corresponding EDS elemental maps of the Stellite 6, 1.0Ti and 1.0Ti20Ni coatings, which are

Fig. 4. SEM images of coatings (a) Stellite 6, (b) 0.5Ti, (c) 1.0Ti, (d) 1.5Ti and (e) 2.0Ti, and (f) the area fraction of precipitates in titanium-modified coatings.

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Fig. 5. SEM images of coatings (a) 1.0Ti10Ni, (b) 1.5Ti10Ni, (c) 2.0Ti10Ni, (e) 1.0Ti20Ni, (f) 1.5Ti20Ni and (g) 2.0Ti20Ni. (d) the area fraction of precipitates in all coatings, and (h) the average diameter of the TiC particles in each modified coating.

representative of the original commercial coating, titanium-modified coatings, and titanium- and nickel-modified coatings prepared in this study, respectively. Elemental segregation was clearly revealed. In the Stellite 6 coating, both chromium and carbon were enriched in the eutectic structure because of carbide rich in chromium. On the contrary, in the 1.0Ti and 1.0Ti20Ni coatings, titanium and carbon were enriched in the carbide particles, whereas chromium segregated to the matrix. No obvious segregation was observed in the distribution of nickel. As γ-Co and ε-Co are allotropic phases, the types of carbides observed in the XRD patterns (Fig. 3) were hard to distinguish based on the SEM and EDS results. EBSD was adopted to characterize the features of the phases in Stellite 6, 0.5Ti, 1.0Ti, 2.0Ti and 1.0Ti20Ni coatings that represented the original commercial coating, titanium-modified coatings and titanium- and nickel-modified coatings respectively, as shown in Fig. 7. The Stellite 6 coating consists of γ-Co matrix, M23C6 and ε-Co. It is noted that M23C6 and ε-Co coexisted, which is indicated by black circle and magnified in Fig. 7a. With the addition of titanium only, TiC was precipitated and ε-Co formed in the γ-Co matrix, as shown in Figs. 7b to d. The eutectic structure consisting of γ-Co, M23C6, and ε-Co (marked with yellow arrows in Fig. 7b) still remained in the 0.5Ti coating. The intermetallic compound in 2.0Ti was still unidentified (marked with yellow arrows in Fig. 7d), though this was subsequently identified through TEM. With the addition of titanium and

nickel, only γ-Co and TiC could be found in the relevant coating (shown in Fig. 7e). No M7C3, which was a phase that could potentially form, based on the prediction of the calculated phase diagram in Fig.1 and the XRD results in Fig. 3, was detected in the coatings mentioned above. In order to identify the types of particles present in the coatings, TEM imaging and selective area diffraction (SAD) were employed. Fig. 8 shows the bright field images of the Stellite 6, 1.0Ti, and 2.0Ti coatings and the SAD patterns (shown in the insets) of the particles corresponding to the white circles for each coating. The EDS results of these particles with numbers are listed in Table 4. Based on the SAD and EDS results, the M23C6 eutectic carbide in the Stellite 6 coating and TiC particles in 1.0Ti were confirmed (Fig. 8a and b). The intermetallic compounds shown in Fig. 8c and d were the same, and the ratio of alloy elements (titanium, chromiun, and iron) to cobalt was 1.0 (Sp3 and Sp4 listed in Table 4). The SAD results in Fig. 8c and d showed that the space group of intermetallic was P4/mnm. The intermetallic was thought to be CrCo. 3.2. Mechanical properties The stress-strain curves of the Stellite 6 and modified coatings, which were obtained from the tensile tests conducted at room

Fig. 6. SEM images and corresponding EDS elemental maps for coatings: (a)–(d) Stellite 6, (e)–(i) 1.0Ti, (j)–(n) 1.0Ti20Ni.

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Fig. 7. Phase maps detected by EBSD: (a) Stellite 6, (b) 0.5Ti, (c) 1.0Ti, (d) 2.0Ti, and (e) 1.0Ti20Ni.

temperature, are plotted in Fig. 9a. All the stress-strain curves showed significant work hardening, followed by sudden fractures without macroscopic necking. Fig. 9b and c summarize the strength and elongation to fracture of all the coatings. Generally, the strengths of all the modified coatings were lower than that of the Stellite 6 coating to some extent. The strength of each modified coating decreased with increasing amount of nickel added. In contrast, the elongation to fracture of most of the modified coatings was higher than that of the Stellite 6 coating. The elongation to fracture was also highly dependent on the nickel concentration; increasing the amount of nickel increased the elongation to fracture. It is noteworthy that the modified coatings whose molar ratio of titanium to carbon was 1.0, namely 1.0Ti, 1.0Ti10Ni and 1.0Ti20Ni, displayed the highest strength and elongation fracture for each nickel concentration. Particularly, the 1.0Ti20Ni coating exhibited the highest elongation of ~12%, with a high ultimate strength of ~1230 MPa. Although its strength only reached ~86% that of Stellite 6, its elongation to fracture was dramatically improved, being four times that of Stellite 6.

Fig. 10 shows the fracture surfaces of the tensile specimens used to examine the plastic deformation behavior of the coatings. A mixed fracture mode was observed in the Stellite 6 coating (Fig. 10a). The tear ridges along the growth directions of dendrites and the eutectic structures were prominent on the fracture surface with the local quasicleavage fracture mode observed on the flat fracture surface. The quasi-cleavage indicated the contribution of brittleness to the fracture process. The tear ridges along specific direction were thought to be associated with the effect of M23C6. As the molar ratio of titanium to carbon increasing, a gradual change could be observed on the fracture surface, as seen in Fig. 10b to d. The tear ridges were partially suppressed, instead presenting small flat quasi-cleavage facets on the fracture surface of the 0.5Ti coating (Fig. 10b). Then, the morphology of fracture surfaces totally transformed to the mode of small flat quasicleavage facets in the cases of the 1.0Ti and 1.5Ti coatings (Fig. 10c and d). Conversely, the fracture surface of coatings with nickel added presented dimples, which were indicative of local ductile fracture (Fig. 10e to g). The 1.0Ti20Ni coatings, which exhibited the highest elongation to fracture, displayed fine and deep dimples on the fracture

Fig. 8. TEM bright field images and SAD pattern of (a) Stellite 6 coating and M23C6 eutectic carbide, (b) 1.0Ti coating and TiC carbide, (c) and (d) 2.0Ti coatings and intermetallic compounds.

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Table 4 Spectrum of EDS test points in Fig. 8 (mol %).

Sp1 Sp2 Sp3 Sp4

C

Ti

Cr

Fe

Co

Ni

W

23.97 35.32 0.89 0.41

/ 61.42 14.03 12.59

62.90 1.04 22.48 22.15

1.03 0.22 11.53 13.65

7.09 0.84 49.24 49.26

0.83 0.03 0.93 0.91

4.18 1.13 0.90 1.03

surface (Fig. 10e). However, the dimples became shallower with increasing molar ratio of titanium to carbon in the titanium- and nickelmodified coatings, as seen on the fracture surfaces of the 1.5Ti20Ni and 2.0Ti20Ni coatings (Fig. 10f and g). The cross-section microstructure of Stellite 6 and the typical modified coatings after tensile tests were characterized by EBSD, and are shown in Fig. 11. The areas observed were close to the fracture sites. The phase map of Stellite 6 after the tensile test, shown in Fig. 11a, indicates the appearance of ε-Co, compared with that of the as-deposited coating, shown in Fig. 7a. The stripped ε-Co phase exhibited a preferred orientation. The map of the boundary distribution corresponding to Fig. 11a, which is illustrated in Fig. 11d, shows a number of angular boundaries between 2° and 10° (marked with blue lines) in the γ-Co matrix. The 2° to 10° angular boundaries displayed a similar preferred orientation as ε-Co. The strip boundaries indicated by red arrow were obviously parallel to the phase boundaries between γ-Co and ε-Co, which are marked with black lines. After the tensile test, the amount of ε-Co phase in the 1.0Ti coating increased, compared with the asdeposited 1.0Ti coating in Fig. 7c. However, unlike the Stellite 6 coating, fewer angular boundaries between 2° and 10° were retained in the residual γ-Co. Instead, phase boundaries were rich in the detected area, as shown in Fig. 11e. In addition, some high angle boundaries (greater than 15°; marked with red lines) could be found in ε-Co. In 1.0Ti20Ni coating, no ε-Co was detected in the phase maps (Fig. 11c). The density of boundaries of the 1.0Ti20Ni coating was lower than that of the other two coatings, as seen in Fig. 11f. Comparing the phase and boundary

distribution maps (Fig. 11c and f), it could be noted that the boundaries angled between 2° and 10° distributed around TiC particles. Few boundaries with angles greater than 10° could be found. Interestingly, twins along a certain direction were detected, which are marked with red lines superimposed on a yellow background in Fig. 11f. Fig. 12 shows the thermal fatigue crack length and life time for Stellite 6 and modified coatings through thermal fatigue crack propagation tests. Only 6 types of modified coatings, that revealed good performances in the tensile tests were subjected to this test: 1.0Ti, 1.0Ti10Ni, 1.0Ti20Ni, 1.5Ti, 1.5Ti10Ni and 1.5Ti20Ni. The number of cycles started to be counted when the cracks length reached 500 μm. Cracks of lengths lower than 10,000 μm were examined in this study. As seen in Fig. 12a, the increase in crack length was almost linear for each coating. The crack propagation rates are summarized in Fig. 12b. All the modified coatings displayed lower propagation rates than the Stellite 6 coating. Nevertheless, the titanium-modified coatings (1.0Ti and 1.5Ti) revealed relatively high mean propagation rates ranging from ~18 to ~22 μm/cycle. However, the titanium- and nickel- modified coatings (1.0Ti10Ni, 1.0Ti20Ni, 1.5Ti10Ni and 1.5Ti20Ni) further decreased the propagation rates. The lowest propagation rate was exhibited by 1.0Ti20Ni, ~4.8 μm/cycle, which was as 6.1% as that of the Stellite 6 coating. Fig. 13 shows the SEM images of the crack tips of the Stellite 6, 1.0Ti and 1.0Ti20Ni coatings after the thermal fatigue crack propagation tests. The modes of crack propagation changed between the Stellite 6 and modified coatings. As seen in Fig. 13a, cracks propagate along the netlike eutectic structure. It is also seen that the eutectic structure debonded with the cobalt matrix, which is marked by a yellow arrow in front of the crack tip. The cracking features of 1.0Ti and 1.0Ti20Ni, shown in Figs. 13b and c, are similar. The crack mainly propagated through the cobalt matrix, apart from propagation along the interface of the particles and the matrix (marked by yellow arrows in Figs. 13b and c). The voids beside the particles are indicted by red arrow in Fig. 13b, which suggests that the interfaces of the particles and the matrix are potential defects that promote crack propagation during thermal fatigue. Nevertheless, the crack sometimes had to propagate through the particle, as indicated by the red arrow in Fig. 13c. The

Fig. 9. Tensile properties of Stellite 6 and modified coatings: (a) stress- strain curve for coatings, (b) strength, (c) elongation to fracture.

Y. Wu et al. / Materials and Design 162 (2019) 271–284

279

Fig. 10. Fracture Surfaces of Stellite 6 and modified coatings: (a) Stellite 6, (b) 0.5Ti, (c) 1.0Ti, (d) 1.5Ti, (e) 1.0Ti20Ni, (f) 1.5Ti20Ni, and (g) 2.0Ti20Ni.

carbide particles, revealing high hardness and strength, also played a role in hindering the propagation of cracks,

4. Discussion 4.1. Effect of titanium and nickel alloying on the solidification microstructures

3.3. SFE of the cobalt-based matrix The SFE of the cobalt-based matrix was calculated through the computational thermodynamic approach, as shown in Fig. 13. The SFE of the Co-Cr-Ni-Ti quaternary cobalt matrix, alloyed with chromium and nickel (0 to 30 mol%) for three different titanium concentration (0 mol%, 2.5 mol% and 5 mol%), decreased for all nickel contents with increasing chromium content, which was consistent with the results of Tria [41]. Nickel would increase the SFE. The range of increase in the SFE as a result of nickel addition depends on the titanium content. It is limited if titanium content is low, while the increase is amplified with increasing titanium content. Titanium by itself significantly improves the SFE even when the nickel content is not increasing. It is important to note that the value of SFE calculated in this study based on the computational thermodynamic method might not be the true value because of the lack of thermodynamic data as described in Section 2.5. Nevertheless, the calculation results could still reveal the effect of the alloying elements on the SFE.

The purpose of adding of titanium was to modify the M23C6 net-like eutectic carbide to obtain isolated TiC particles. Titanium is a strong carbide-forming element [43]. In the melted pool, the carbon atoms first interact with the titanium atoms to form TiC, and then, the remaining carbon interact with the chromium atoms to form the chromiumrich M23C6, because the that chemical affinity between carbon and titanium is stronger than that between carbon and chromium [44]. As shown in Fig. 4a to e and Fig. 7a to d, isolated TiC precipitated in the cobalt matrix that gradually replaced the eutectic carbide as the titanium content increased. Referring to the calculated phase diagram shown in Fig. 1a, when the concentration of titanium is less than ~4.70 wt%, a eutectic structure appears during the solidification process at the eutectic reaction point. The 0.5Ti coating, whose titanium concentration was 2.14 wt%, revealed this eutectic structure. However, the eutectic carbide was M23C6, identified by EBSD (Fig. 7b) instead of M7C3 predicted by phase diagram (Fig. 1a). Although it is noted that M7C3 might change to M23C6 during the subsequent cooling process, as shown in the

Fig. 11. EBSD maps of Stellit 6 and modified coatings after tensile tests: (a–c) phase maps and (d–f) boundary distribution maps of (a and d) Stellite 6 coating, (b and e) 1.0Ti coating, and (c and f) 1.0Ti20Ni coating.

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Fig. 12. Thermal fatigue crack growth for chosen modified coatings: (a) curves of thermal fatigue crack length versus thermal cycle and (b) mean growth rate.

phase diagram, the occurrence of this phase transformation was difficult during the laser deposition process. The M7C3 → M23C6 transformation is based on metallic atoms (chromium and molybdenum) diffusion and the reaction of 9 M + 2M7C3 → M23C6. It is hard to reach equilibrium situation under the ultrahigh cooling speed of 104–106 K/s in the laser deposition process. According to Wieczerzak et al. [45] and Inoue et al. [46], increasing the chromium/carbon ratio in the alloy resulted in the preferential formation of M23C6, rather than M7C3, in the nonequilibrium state, especially in a rapid solidification process. It is assumed that M23C6 formed directly through the eutectic reaction, rather than through the transformation from M7C3. With increasing titanium concentration, the 1.0Ti and 1.5Ti coatings revealed only the primary MC carbide (TiC) in the cobalt matrix, because their solidification corresponded to the γ + MC phase region. It is noteworthy that the 1.0Ti coating only contained 4.20 wt% titanium content, which was less than 4.70 wt%. The eutectic structure would be theoretically formed in the equilibrium solidification state in the 1.0Ti coating. However, the high cooling speed of the melted pool, combined with sufficient titanium content (molar ratio of titanium to carbon was 1.0) that was prone to formed MC carbide with carbon, ensured that the solidification process of the 1.0Ti coating occurred in the γ + MC phase region. The increase amount of titanium was limited because of the formation of intermetallic compounds in the 2.0Ti coating. The intermetallic compounds deteriorated the properties of the coating. One interesting phenomenon that was observed was that the chromium in the modified coatings was mainly distributed in the cobalt matrix rather than being rich in carbides, compared to the Stellite 6 coating, as seen in Fig 6c, g, and l. Chromium is a strong ε-Co phase stabilizing element [35,47]. As seen in Fig. 14, increasing the chromium content significantly decreases the SFE for all nickel and titanium contents, which resulted in the cobalt-based matrix readily undergoing the athermal martensitic transformation during cooling [48,49]. The ε-Co phase was detected in the matrix of a titanium-modified coating, as shown in Fig. 7b to d. The addition of nickel was aimed at improving the ductility of the cobalt matrix by stabilizing the γ-Co phase. Based on the computational results presented in Fig. 14, the SFE improved with increasing nickel content, which suggested that the γ-Co phase tended to exist at room temperature. The appearance of the ε-Co phase, which is detrimental to ductility [50], in the titanium-modified coatings was not expected.

Therefore, when plastically forming the alloys, a large amount of nickel (10 to 37 wt%) is generally added to Co-Cr-Mo or Co-Cr-W alloys to improve the ductility [48]. In this study, the addition of nickel (10 and 20 wt%) to the titanium-modified coatings restrained the formation of the ε-Co phase effectively, which was confirmed by the XRD and EBSD results shown in Fig. 3 and 7e, respectively. Only γ-Co phase could be found in the nickel-modified coatings. It is noted that combined addition of titanium and nickel to a cobalt-based matrix is more efficient in increasing the SFE than the addition of only nickel, as seen in the SFE results presented in Fig. 14. 4.2. Effect of carbide morphology on the mechanical properties The addition of titanium resulted in the M23C6 net-like eutectic carbide changing to the fine isolated TiC, as seen in Fig. 4a to d. The elongations to fracture of the titanium-modified coatings slightly increased, but the ultimate tensile strength and 0.2% proof stress dropped to some extent (Fig. 8b and c). The 1.0Ti coating exhibited the best strength and ductility among all the titanium-modified coatings. Interestingly, the area fraction of the carbide precipitate of the 1.0Ti coating was the smallest (Fig. 4f). All these results indicated that the morphology and content of carbides were the factors that influenced the strength and ductility of the modified coatings. The directional dendrite structure characteristics and tear ridges along the dendrite growth direction were revealed on the fracture surface of Stellite 6 (Fig. 10a), indicating the weakness at the dendrite interface. These fracture features were attributed to the net-like eutectic carbides in the interdendritic region, and these carbides were expected to trigger fractures during tensile deformation [51,52]. The void was prone to nucleate at the interface between eutectic carbide and the matrix, and usually developed into relatively large size [53], inducing premature fracture. These factors account for the low ductility of the Stellite 6 coating. In contrast, fine and dispersed carbides in the modified coatings facilitated the deformation of the cobalt-based matrix and had a pinning effect on the dislocations that helped in retaining the strength of the coatings [54,55]. As seen in Fig. 11f, a large number of low-angle boundaries (between 2° and 10°) appeared around the TiC particles, which indicated that the deformations accumulated. This phenomenon accounted for the pinning effects of the TiC particles. However, the elongations to

Fig. 13. SEM images of the of crack tips of Stellite 6 and modified coatings after thermal fatigue: (a) Stellite 6coating, (b) 1.0Ti coating, and (c) 1.0Ti20Ni coating.

Y. Wu et al. / Materials and Design 162 (2019) 271–284

281

Fig. 14. Calculated SFE of cobalt-based matrix: (a) Co-xCr-yNi-0Ti system, (b) Co-xCr-yNi-2.5Ti system, and (c) Co-xCr-yNi-5Ti system.

fracture of the titanium-modified coatings were still relatively low, between 4% and 5.7% (Fig. 9c). As mentioned above, the addition of titanium led to the appearance of ε-Co phase in the matrix (Fig. 7b to d). It is well known that the ε-Co phase has poor plastic deformability because of its limited slip systems [56]. This accounts for the low ductility of titanium-modified coatings. 4.3. Effect of SFE on the mechanical properties The deformation capability of the cobalt-based matrix played an important role in improving the ductility. Mori et al. [56] pointed out that it was important to retain the γ-Co single phase in the initial microstructure and suppress the γ/ε martensitic transformation during deformation. Both these factors were greatly affected by the SFE [56–58]. As seen in Fig. 11d to f, the deformation models of three typical coatings (Stellite 6, 1.0Ti and 1.0Ti20Ni) with different SFEs presented distinguishing features. According to the computational results (Fig. 14), the SFE values of the three coatings followed the order 1.0Ti ﹤ Stellite 6 ﹤ 1.0Ti20Ni. Comparing the EBSD results of the coatings after the tensile tests (Fig. 11) with those obtained in the as-deposited state (Fig. 7a, c and e), it is noted that the ε-martensitic transformation was more likely to occur during the tensile tests in the 1.0Ti coating. The γ-Co in the 1.0Ti coating represented a metastable phase at room temperature that was readily converted to ε-Co through strain-induced martensitic transformation during the tensile tests [50,59]. The ε-Co martensite formed and crossed the already existing plate-like athermal martensitic ε-Co phase; early fracture might result due to the stress concentration generated [56]. The interface of γ-Co and ε-Co could become the crack propagation route [59]. As shown in Fig. 10c the fracture morphology of the 1.0Ti coating presented a quasi-cleavage fracture with small facet-like features. This is because the γ-Co phase is fragmented into small pieces by the strain-induced ε-Co martensite (Fig. 11b), which represented the crack initiation site and the crack propagation route. The 1.0Ti20Ni coating had a high SFE. No ε-Co martensite was detected in the as-deposited state (Fig. 7e) and after the tensile test (Fig. 11c). The γ-Co matrix of the 1.0Ti20Ni coating was more stability during the deformation. Interestingly, a few strip-like twins were observed in the deformed cobalt matrix. Rémy [60] pointed out that the deformation twinning occurred concurrently with the propagation of undissociated dislocations for increasing amounts of deformation when the SFE is increased. Excellent deformation capability and no potential defect sites (like ε-Co martensite) increased the elongation to fracture, resulting in a dimple-like fracture morphology being observed in the 1.0Ti20Ni coating (Fig. 10e).

shown in Fig. 5d, with the addition of titanium, the fraction of carbides dropped from ~25% in the Stellite 6 coating to ~16% and ~19% in the 1.0Ti and 1.5Ti coatings, respectively. Correspondingly, the mean thermal fatigue crack propagation rates of the 1.0Ti and 1.5Ti coatings decreased to ~21 μm/cycle and ~18 μm/cycle, respectively, from ~79 μm/ cycle for Stellite 6 (Fig. 12b). In other words, a 24–36% decrease in the fraction of carbides resulted in 73–79% dramatic reduction in the crack propagation rate. The isolated TiC particles formed could effectively suppress the thermal fatigue crack propagation, compared to the netlike eutectic carbides. The addition of nickel decreased the amount of carbides and the thermal fatigue crack propagation rate. The relationship between the decrease in the amount of carbides and the decrease in the thermal fatigue crack propagation rate is shown in Fig. 15. A positive linear correlation was observed for the two series of coatings based on different molar ratios of carbon to titanium. One series corresponded to the molar ratio of 1.0: 1.0Ti, 1.0Ti10Ni, and 1.0Ti20Ni. The other series corresponded to the molar ratio of 1.5: 1.5Ti, 1.5Ti10Ni, and 1.5Ti20Ni. It is noted that the slopes of the two regression lines were less than 1, which indicated that a decrease in the amount of carbides might not result in an equivalent reduction in the thermal fatigue crack propagation rate. Therefore, modifying the morphology of the carbides was more effective in reducing the thermal fatigue crack propagation rate than decreasing the amount of carbides. The addition of nickel, on one hand, could reduce the amount of carbides. On the other hand, the addition of nickel stabilized the γ-Co phase in the matrix and increased the deformation capability. It is known that the thermal fatigue process resulted in materials suffering

4.4. Effect of carbides and SFE on the thermal fatigue crack propagation rates The thermal fatigue crack propagation rate was mainly affected by the morphology of carbides rather than by the amount of carbides. As

Fig. 15. The relationship between the decrease in amount of carbides versus the decrease in thermal fatigue cracks propagation rate.

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from plastic strains to some extent. In addition, the area near the crack tip was subjected to a more serious plastic strain because of the stress concentration there. Excellent deformability helped to accommodate more dislocations. The γ-Co matrix stabilized by nickel revealed good plasticity, as shown by the tensile tests. Therefore, the thermal fatigue cracks were suppressed when they grew into the γ-Co matrix.

displayed the maximum elongation to fracture (~12%), along with a relatively high strength (1230 MPa). 4. The isolated TiC particles and γ-Co matrix could effectively suppress the thermal fatigue crack propagation. A linear correlation was observed between the decrease in the amount of carbides and the decrease in the fatigue crack propagation rate. The 1.0Ti20Ni coating presented the lowest thermal fatigue crack propagation rate.

5. Conclusions Acknowledgements In this study, titanium and nickel were added to Stellite 6 commercialized coatings prepared by laser deposition to modify the microstructure and the properties. The concept was verified based on the amounts of titanium and nickel added. The molar ratio of titanium to carbon in modified coatings was 0.5, 1.0, 1.5, and 2.0. Nickel was further added in 10 and 20 wt% to the titanium-modified coatings. The microstructures, tensile properties, and thermal fatigue crack propagation rates of the modified coatings were examined. The following conclusions were drawn: 1. The addition of titanium modified the morphology of the carbides from the M23C6 net-like eutectic carbide observed in the Stellite 6 coating to isolated TiC. Besides, the reaction of titanium with carbon resulted in the chromium originally contained in M23C6 segregating in the cobalt matrix, which decreased the SFE and induced the appearance of ε-Co phase. 2. The addition of nickel stabilized the γ-Co phase by increasing the SFE. 3. The strengths of the titanium- and nickel-modified coatings slightly decreased, but the elongations to fracture increased, compared with the Stellite 6 coating. These were attributed to the dispersed isolated TiC particles and the stable γ-Co phase. The 1.0Ti20Ni coating

This work was supported by the National Natural Science Foundation of China (grant numbers 51474178) and National Key and Program of China (grant numbers 2016YFB1100202). We thank Li Yang in Chong qing University for her help in phase diagram calculation. Appendix A. Computational thermodynamic approach The formulation to calculate the intrinsic stacking faults energy (SFE) as proposed by Olson and Cohen [39,41] is: γsf ¼ 2ρΔGγ→ε þ 2σ γ=ε

ðA1Þ

The detailed calculation method can be referred to the works of Achmad [41], as shown in Fig. A1. The meaning of each parameter is listed in Table A1. The thermodynamic data required for the calculations and their references was list in Table A2. The data required for calculate the magnetic contribution (ΔGmgγ→ε) and their references are listed in Table A3.

Fig. A1. The flowchart of intrinsic SFE calculations by the computational thermodynamic approach provided by Achmad [41].

Y. Wu et al. / Materials and Design 162 (2019) 271–284 Table A1 Parameter used in SFE calculation. Parameter Meaning γsf ΔGγ→ε σγ/ε a E

Gi, j, γ→ε k

The intrinsic SFE The Gibbs free energy difference of γ-fcc to ε-hcp phase transformation The interfacial energy per unit area of the phase boundary Lattice parameter Ternary excess Gibbs free energy

Unit

Value

mJ/m2

/

mJ

/

mJ/m2

7.5

Å 3.544 mJ/m2 /

Table A2 of the fcc Thermodynamic function describing the change in the Gibbs free energy ΔGγ→ε i to hcp phase transformation for the pure elements, the excess free energy coefficients Ωijγ→ε γ→ε E and the ternary excess Gibbs free energy Gi,j,k used in this study. Parameter Thermodynamic function (J/mol) ΔGCoγ→ε ΔGCrγ→ε ΔGNiγ→ε ΔGTiγ→ε ΩCoCrγ→ε ΩCoNiγ→ε ΩCoTiγ→ε ΩCrNiγ→ε ΩNiTiγ→ε E GCoCrNiγ→ε E GCoCrTiγ→ε

−427.59 + 0.615 T −2846 − 0.163 T 1046 + 1.255 T −6000 + 0.1 T −4621.59 + 7.32 T + (7341.73–7.93 T)(XCo − XCr) −820 − 1.65 T 65,258 − 29.109 T-16240(XCo − XTi) 27 + 12.1 T + (−19,895 + 16.38 T) (XCr − XNi) 118,143 − 6.706 T + 62,430(XNi − XTi) XCo XCr XNi (−28 T XCr) XCo XCr XTi ((−324,295 + 45 T)XCo + (137,386 + 45 T) XCr)

Reference [42] [42] [42] [42] [61] [62] [63] [62] [64] [62] [65]

Table A3 The formulas for the magnetic moment of Φ phase (βΦ) and the Curie temperature for magnetic ordering of the Φ structure (TΦ c ) (K) used in this study. System Function Tc, Coε−hcp = 1396 Tc, Coγ−fcc = 1396 βCoε−hcp = 1.35 βCoγ−fcc = 1.35 Pure Cr Tc−equ, Crε−hcp = − 1109 Tc−equ, Crγ−fcc = − 1109 β−equ, Coε−hcp = − 2.46 β−equ, Coγ−fcc = − 2.46 Pure Tc, Niε−hcp = 633 Ni Tc, Niγ−fcc = 633 βNiε−hcp = 0.52 βNiγ−fcc = 0.52 Co-Cr Tc, Co−Crε−hcp = −5828.68 XCo XCr + 4873.95 (XCo − XCr) XCo XCr Tc, Co−Crγ−fcc = −9392.53 XCo XCr + 8383.04 (XCo − XCr) XCo XCr Co-Ni Tc, Co−Niε−hcp = 411 XCo XNi − 99 (XCo − XNi) XCo XNi Tc, Co−Niγ−fcc = 411 XCo XNi − 99 (XCo − XNi) XCo XNi βCo−Niε−hcp = 1.046 XCo XNi + 0.165 (XCo − XNi) XCo XNi βCo−Niγ−fcc = 1.046 XCo XNi + 0.165 (XCo − XNi) XCo XNi Pure Co

Reference [42] [42] [28,42,61] [28,42,61] [42,66] [42,66] [61,62]

[42,62] [42,62]

Noted that: the pure Cr do not possesses Tc and β values of fcc and hcp phases, the equivalent effects on these two values was calculated according to references [28, 42, 61].

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