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TiAl composite coating prepared by laser cladding
Optics and Laser Technology 112 (2019) 339–348
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Full length article
Effects of TiC on the microstructure and properties of TiC/TiAl composite coating prepared by laser cladding
T
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X. Hea,b,c, R.G. Songa,b,c, , D.J. Kongb,d a
School of Materials Science and Engineering, Changzhou University, Changzhou, Jiangsu 213164, China Jiangsu Key Laboratory of Materials Surface Science and Technology, Changzhou University, Changzhou, Jiangsu 213164, China c Jiangsu Collaborative Innovation Center of Photovolatic Science and Engineering, Changzhou University, Changzhou, Jiangsu 213164, China d School of Mechanical Engineering, Changzhou University, Changzhou, Jiangsu 213164, China b
H I GH L IG H T S
prepared TiC/TiAl composite coatings on TiAl alloy by laser cladding. • We microstructure of coating containing nano-TiC was finer and less defects. • The coating can improve the hardness and wear resistance of the substrate. • TiC/TiAl • The coating can significantly improve the corrosion resistance of the substrate.
A R T I C LE I N FO
A B S T R A C T
Keywords: TiAl alloy Laser cladding TiC/ TiAl composite coating Microstructure Properties
In order to study the effect of TiC on the microstructure and properties of cladding coatings, TiC/ TiAl composite coatings have been prepared by using a laser cladding technique on TiAl alloy, and the microstructure and properties of the resulting composite coatings have been investigated using scanning electron microscopy (SEM), X-ray diffraction (XRD), micro-hardness testing, X-ray stress measurements, friction and wear testing, and an electrochemical workstation. The results showed that when the particle size of TiC is up to micron, the growth of TiC is more developed, the dendrite growth direction is disordered, and the cladding quality is poor. When the particle size of TiC is nanoscale, the morphology of the reinforcement phase TiC is mainly granular and thin rods, which is uniform in the coating and the direction of growth is regular. When the concentration of TiC reaches 20%, the microstructure of the cladding layer grows well and is dense. The microhardness and wear resistance of the coatings prepared by adding nano TiC are better than those of the coatings prepared by micron TiC. When the content of nano TiC in the coatings increases from 10% to 20%, the microhardness and wear resistance of the coatings are obviously improved. The residual stress of the coating is positively related to the crack rate. The residual stress in the coating is tensile stress, and the crack type is the crystal crack in the hot crack. When the particle size of TiC is small and the concentration is high, the coating shows better corrosion resistance.
1. Introduction TiAl alloy is a new kind of intermetallic compound structure material, which has many outstanding characteristics, such as low density, high specific strength and specific modulus, high specific rigidity and good high-temperature creep resistance. The combination of these advantages will push TiAl alloy to be used in automotive, aerospace and marine industries. Therefore, TiAl intermetallic compounds have attracted wide attention. However, TiAl alloys display poor high-temperature oxidation resistance as high-temperature parts. Moreover, the
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low hardness and poor tribological properties severely limit their applications as key moving components. When exposed to the sea water environment, they are not immune to combined corrosion and wear attacks. Although protective passive films with a thickness of 3–5 μm readily form on the surface of titanium alloy in sea water, they are too thin to resist the shear stress and removed quickly, accelerating the metal dissolution and corrosion. In order to overcome this drawback, the formations of ceramic coatings on the TiAl surface are proposed and proven to be the effective methods to improve the surface properties of alloys. Among various materials, the TiAl intermetallic based coating is
Corresponding author at: School of Materials Science and Engineering, Changzhou University, Changzhou, Jiangsu 213164, China. E-mail address: [email protected] (R.G. Song).
https://doi.org/10.1016/j.optlastec.2018.11.037 Received 4 July 2018; Received in revised form 21 October 2018; Accepted 18 November 2018 0030-3992/ © 2018 Elsevier Ltd. All rights reserved.
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residual stress were as follows: fixed peak method for cross correlation, using a Co target, Co-Kα1 radiation source, Prague crystal face of (4 0 0), incident angles of 0°, 25°, 35°, and 45°, stress constant of −130 MPa/°, 2θ scan start and termination angles of 155–145°, 2θ scanning step of 0.10°, counting time of 0.50 s, X-ray light tube high voltage of 22 kV, and an X-ray light tube current of 6.0 mA. Electrochemical tests were carried out by using CS350 type electrochemical workstation. The sample dimension of electrochemical test was 10 mm × 10 mm × 3 mm. Except for the surface of the coating, the rests of the surfaces were inlaid with epoxy resin. The test medium was a 3.5% NaCl solution. The system of electrochemical workstation included a saturate calomel electrode (SCE) as a reference electrode, Pt as a counter electrode, sample as working electrodes. The potentiodynamic scanning rate of 1 mV/s, sampling frequency of 0.5 Hz, the test potential range of −0.5 to 0.5 V, and test time of 1800 s. EIS measurement's frequency range is 10−1 to 105, and test time of 300 s. All the measurements were performed after 30 min immersion in 3.5% NaCl solution until open circuit potential (OCP) was stable.
a promising method to enhance the high-temperature oxidation resistance, corrosion and tribological properties of TiAl alloys [1–3]. In recent years, in order to further study TiAl alloys, many researchers have improved the wear resistance, high-temperature oxidation resistance and corrosion resistance of TiAl intermetamerics by a series of surface modification techniques such as ion infiltration, deposition, laser alloying, and laser cladding [4–7]. Laser cladding is an advanced surface modification technology. The thickness of the cladding layer obtained by laser cladding is large, and the structure is dense and uniform [8–9]. Reinforced carbides TiC possess both excellent hightemperature abrasive and corrosion resistance because of their very high hardness, unique strong atomic bonds and high stability under high-temperature exposure. So the high-temperature coatings reinforced by TiC carbides are expected to possess good high-temperature corrosive wear resistance [10]. At present, Chai et al. studied the effect of nano-TiC/C on the microstructure and properties of laser cladding Ni-base alloy coating. The results show that with the increase of nanoTiC/C addition, the number of γ-Ni dendrites showed a trend of increasing first and then decreasing, and the change trend of its size was reversed [11]. Liu has also studied the preparation of γ/W2C/TiC composite coatings on TiAl alloys. The results show that the wear resistance and high-temperature oxidation resistance of the coatings are significantly improved [12]. At present, there are few studies at home and abroad to study the effect of particle size and concentration of TiC on the microstructure and properties of laser cladding composite coatings [13–15]. Therefore, this paper adopts laser cladding technology to prepare TiC-reinforced composite coatings on TiAl alloys, and analyzes the influence of TiC on the structure, hardness, wear resistance and corrosion resistance of cladding layers.
3. Results and discussion 3.1. XRD and morphologies analysis of powders Fig. 1 shows the X-ray diffraction pattern of the mixed powder used for the laser cladding test. Because of the lower milling speed, the mixed powder did not produce a new phase during ball milling, and it was still Ti, Al and TiC phases [16]. Fig. 2 shows the micromorphology of TiC powder at nanometer and micrometer diameters. The nano TiC used in the experiment is spherical and has a particle size of about 50 nm. The micro TiC used is in the form of a block or a sheet with irregular morphology. The particle size is about 40 μm, which contains fine TiC debris.
2. Experimental The experimental material and contrast material was TiAl alloys, which are mainly composed of Ti-48Al-2Cr-2Nb (at.%). The sample size is φ46 × 9 mm. Before the experiment, the specimens were sanded step by step with sandpaper and washed repeatedly with anhydrous ethanol. The spherical powders used in the cladding test are mainly Ti powder (purity 99.3%, 64 μm–120 μm) and Al powder (purity 99.0%, 220 μm–330 μm). Ti powder and Al powder are mixed according to the atomic ratio of 1:1, then three different TiC powders were added. 20% micron TiC was added to the group A powder, 10% nano-TiC were added to the group B powder, and 20% nano-TiC were added to the group C powder. The coating was chemically pre-prepared, and methylcellulose was used as a binder. The thickness was about 1 mm, and it was then incubated in a drying oven at 80 °C. for 2 h. The laser power was fixed at 2.40 KW, the scanning speed was 10 mm/s, the multi-pass rate is 40%, the laser cladding adopts a circular spot, the diameter is 3 mm, the above process parameters are fixed, and the preparation contains 20% TiC (μm) A coating, 10% TiC (nm) B coating, 20% TiC (nm) C coating. The morphologies, element compositions and phases of coatings were analyzed using a JSM-6360LA type scanning electron microscope (SEM), energy dispersive spectrometer (EDS), and D/max2500 PC X-ray diffraction (XRD), respectively. The microhardness of the coating was measured with an THL200 digital microhardness tester with the following conditions: a loading of 200 g and a loading time of 15 s. The hardness of cladding layer is measured every 100 μm from the cladding surface to the substrate, and the average value is obtained three times at the same depth level. Friction and wear testing were conducted using a CFT-1 surface tester. The grinding material used was SiC at a loading of 150 N and a motor speed of 400 r/min, using a reciprocating sliding mode, the wear scar radius was 3 mm and the running time was 30 min. The wear medium was air and the measurement was completed using a BT25S electronic analytical balance to measure the weight loss. The residual stresses of the coatings were analyzed using an X-ray stress tester. The technological parameters used for measurement of the
3.2. Morphologies of coating interfaces Fig. 3 shows the metallographic cross-section of the coating. It can be seen that the cross-section of the cladding layer is in turn divided into four zones: the coating zone (CZ), the diffusion zone (DZ), the heat affected zone (HAZ) and the substrate. In the A coating, the microstructure of the cladding layer is dense, and the thickness of the cladding layer is about 0.8 mm. The reinforcing phase TiC is scattered in the cladding layer and has no obvious gradient distribution characteristics [17]. In the B coating, the structure of the cladding layer is sparse, the thickness is about 0.75 mm, and the reinforcing phase TiC is more regularly distributed in the coating, and it can be seen that the
Fig. 1. XRD patterns of the mixed powder. 340
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(b)
(a)
100 nm
100 ȝm
(c)
(d)
Fig. 2. Morphology of TiC with different particle sizes and the diameter distribution of the TiC powder. (a, c) nano TiC (b, d) micron TiC.
considered that the dilution rate of the coating is better at 5%.
transition between the cladding layer, and the substrate is natural. The microstructure of C coating was well-developed and dense. The thickness of cladding layer was about 0.7 mm. The microstructure of the coating grows up regularly, and the TiC distribution of the reinforcing phase is more uniform. Because the dilution rate directly affects the performance of the coating, the dilution rate λ is expressed by the formula
λ=
h × 100% H+h
3.3. XRD analysis of coating surfaces Fig. 4 shows the XRD pattern of the surface of the coatings obtained by using A, B and C three groups of powders at P = 2.40 kW and V = 10 mm/s. It can be seen that the major components of the A coating are For TiAl and TiC phases, compared with B and C coatings, Ti3Al does not appear at about 40° for A coating, and the main composition phase for B and C coatings is TiAl and TiC phases and a small amount of Ti3Al phase.
(1)
where H is the coating thickness and h is the melting depth of substrate. The melting depth of A, B and C coatings is 0.082 mm, 0.052 mm, 0.039 mm respectively. So the dilution rates of coatings A, B and C can be calculated as 9.03%, 6.51% and 5.18%, respectively. It is generally
Fig. 3. Micrographs of coating cross-section (a) A coating (b) B coating (c) C coating. 341
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d, the main elements in the cladding layer are Ti, Al and C. The atomic ratio of Ti to Al is about 2:1. Therefore, in addition to the TiAl phase, the cladding layer should also contain Ti3Al phases. This is consistent with the results of XRD pattern analysis in Fig. 4. Therefore, the main composition phase in the coating is the continuous phase TiAl. TiC is the reinforcing phase and also contains a small amount of Ti3Al phase. Due to the small TiC content in B powder, the growth phase of TiC grows slowly, and at the same time. The nano TiC in B powder has a spherical shape and regular morphology, resulting in a more uniform distribution in the cladding layer and a more regular growth direction. There are some tiny pores on the surface of the coating, no cracks appear, and the cladding quality is good. Fig. 5c shows the typical microstructure of C coating. Combined with the results of energy spectrum analysis of Fig. 6e and f, the main elements in the cladding layer are Ti, Al, C, and the atomic ratio of Ti to Al is about 2:1, The continuous phase is TiAl phase and a small amount of Ti3Al phase in the C coting. The reinforcing phase TiC is mainly in the form of particles, fine rods. Granular TiC gathers next to the thin rod-shaped TiC, grows according to a certain rule, the dendrite arms in the cladding layer are finer. The organization grows well. Compared with the B coating, the TiC mass fraction in the C coating is 10% greater than that of the B coating, resulting in more dense TiC in the C coating. The surface pores of the coating are further reduced, so that the quality of the cladding layer is improved. Combined with the above test analysis results, it can be found that the coating quality of the A coating is poor, and the coating quality of the B and C coatings is better. Because the microstructure of the crystal has a decisive influence on its growth morphology and mode [19], when the TiC particle size in the A coating and the B and C coatings is greatly different, the A coating is exposed to a high-energy laser beam. Micron TiC absorbs less energy than nano-TiC in the other two coatings, causing TiC in A coating to not fully melt, affecting the temperature field and element distribution in the molten pool, making TiC provide a nucleation surface for crystals in the molten pool, further causing heat stress concentration, the formation of cracks and holes. Comparing B and C coatings, the microstructure and growth direction are not much different. Granular TiC and thin rod-shaped TiC are connected with each other according to certain rules. However, the reinforcement phase TiC is denser and has fewer micropores on the surface of the C coating, which is undoubtedly due to TiC with more C coatings.
Fig. 4. XRD patterns of cladding layer.
3.4. Morphologies and EDS analysis of coating surfaces Figs. 5 and 6 shows the typical microstructures and their energy spectrum analysis results for A, B and C coatings, respectively. It can be seen from Fig. 5a that the reinforcement phase TiC growth of the A coating is relatively well-developed, the structure of the coating is larger and the direction of dendrite growth is disorderly. It is not typical. At the same time, it can be seen that there are a large number of tiny pores and a small number of cracks on the surface of the cladding layer, the coating quality is poor. On the one hand, it may be due to the irregular shape of TiC in the powder A and the small fines that cause stress concentration and cracks on the top of the pores. On the other hand, the thermal physics coefficient between micro TiC and the matrix is very different. This can easily lead to the generation of thermal stress, resulting in cracks and pores, making the cladding less effective [18]. According to the energy spectrum analysis of Fig. 6a and 6b, the main elements in the coating are Ti, Al and C. The ratio of Ti and Al atoms are close to 1:1. Therefore, the continuous phase (Spectrum 1) in the cladding layer is mainly TiAl. Which is mainly due to irregular shape of TiC in A powder, poor fluidity, large particle size, and more laser energy required for cladding, thus affecting its temperature field and element field distribution in the molten pool, making they do not diffuse to each other at the bottom of the bath and do not reach the temperature needed to form the liquid phase, and thus the Ti3Al phase cannot be formed. The reinforced phase in the cladding layer is TiC (Spectrum 2), which is consistent with the XRD pattern analysis results in Fig. 4. Fig. 5b shows a typical tissue morphology of the B coating. The dendrites in the cladding layer are mostly connected in a granular or elongated form. According to energy spectrum analysis of Fig. 6c and
3.5. Microhardness analysis The microhardness distribution from the cladding surface to the substrate is shown on Fig. 7. The microhardness of C coating increases slowly from 0 to 500 μm, and the microhardness reaches a maximum of 813.3 HV0.2 at a distance of about 500 μm from the surface, then the microhardness gradually decreases during the extension to the
(b)
(a)
(c)
Spectrum1
Spectrum 5
Spectrum 3
Crack Spectrum 2
Pore
50 ȝm
Spectrum 6
Spectrum 4
50 ȝm
50 ȝm
Fig. 5. The surface SEM images of the TiC/TiAl coating (a) A coating (b) B coating (c) C coating. 342
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Fig. 6. EDS spectrum of cladding layer of the tested TiC/TiAl coating (a, b) A coating (c, d) B coating (e, f) C coating.
microhardness of the substrate is about 430 HV0.2, and the average microhardness of the A, B, and C coatings are 701.51 HV0.2, 627.86 HV0.2, and 598.04 HV0.2, respectively. Compared with the substrate, the coating has a significantly increased microhardness, especially for the C coating, the microhardness is increased by nearly 50% due to the rich reinforcing phase TiC in the coating surface. However, it can be clearly observed that the microhardness of the B and C coatings is better than that of the A coating. On the other hand, since the B and C coatings are nano reinforcing phase TiC, the resulting microstructure is finer and the fine-grain strengthening effect is more pronounced. In addition, B and C coatings contain Ti3Al phase, which has close-packed hexagonal superlattice structure, high specific strength and high elastic modulus. which can significantly improve the surface hardness of the cladding layer [20].
3.6. Friction and wear testing Fig.7. Microhardness distribution of the TiC/TiAl coatings.
Fig. 8 shows the data of the wear loss tests for the three composite coatings and substrates, it can be seen that the wear of the A, B and C coatings are decreased in varying degrees compared to the original TiAl alloy, the wear loss of the C coating is significantly better than the A and B coatings. On the one hand, due to the distribution of a large number of nanoparticle reinforced phases TiC in the C coating, it has a
substrate. The microhardness of B coating gradually increases from 0 to 300 μm, and reaches a maximum of 798.3HV0.2 at a distance of about 300 μm. The microhardness of A coating reaches a maximum of 669.1 HV0.2 at 300 μm, and then gradually decreases. The average 343
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coatings are 1.14, 1.17, and 1.21, respectively. Therefore, the hardness and wear resistance of the cladding layer can be greatly improved when the TiC concentration of is higher and the grain size is smaller [22]. Fig. 10 demonstrates the surface morphologies of the coatings and substrate afte wearing test. Fig. 10a shows the wear profile of the A coating. It can be seen that the worn surface is rough and distributed with more grooves. The surface has a layer of shallow furrows, and the wear mechanism is mainly micro-cutting and scratching. Fig. 10b shows the wear profile of the B coating. It can be observed that the surface of the coating is relatively flat, with some hard particles on the surface. There are some tiny pits and scratches, but it is not obvious. This is mainly due to the hardness of the B coating. Higher, the cutting effect on the coating weakens. On the other hand, the hard reinforcing phase TiC protrudes on the surface of the substrate and reduces the contact surface with the grinding wheel, but as the wear progresses, the protruding TiC will fall off from the coating surface, and thus Leave pits. The wear surface of the C coating is smoother than that of the B coating, with only minor scratches and partial damage. The wear surface is basically free of plastic deformation and adhesive wear. As the high hardness of nano-TiC diffuses on the surface of the coating, microcutting during grinding is further weakened, and therefore only slight scratches are generated. At the same time. The presence of high hardness TiC makes it play a role in the hard skeleton support of the cladding layer, thereby effectively reducing the wear rate [23]. The wear surface of the substrate is very uneven as shown in Fig. 10d. After the wear, the furrow is deep and wide, there are many traces of friction, the direction is disordered, and there are shallow peels locally. The wear mechanism is mainly abrasive wear.
Fig. 8. Total wear loss of the TiC/TiAl coatings.
higher hardness and strength, and it plays a role in abrasion against in the process of rubbing with the friction pair and hinders the plastic deformation of the coated substrate, and helps to improve the wear resistance of the coating. On the other hand, when the molten pool solidifies at high speed, some elements such as Cr and Nb were infiltrated into the substrate, which played a solid solution strengthening role in the coating. Meanwhile, the dense structure and lower cracks and pores in the B and C coatings are also important factors to enhance the wear resistance of the coating [21]. With the passage of time, the wear loss appears to be gradually increasing. This is due to the high brittleness of TiAl alloys at room temperature. The shedding of TiC inclusions in the wear process forms a new abrasive between the two pairs of grinding parts, and TiC has high hardness, which can be cut into the cladding layer and make the wear increase. Fig. 9 shows the air friction coefficient as a function of time for coatings and substrate, it can be seen from Fig. 9 that the friction coefficient increases rapidly to about 0.6 in 0–3 min, and finally the friction coefficient stabilizes at about 0.5. Compared with the A coating, due to the higher hardness of the B and C coatings, the friction coefficient is further reduced, and finally stabilized at 0.45 and 0.35, respectively. The relative wear resistance is one of the important criteria for judging frictional wear performance, and the relative wear resistance (ε) is defined as
ε=
W0 W1
3.7. Residual stress analysis The cracks were related with residual stresses in the Laser cladding test, from the 2θ-sin2φ regression curves, the residual stress of Cladding coating A, B and C was 143.3 MPa (Fig. 11a), 69.7 MPa (Fig. 11b), 54.1 MPa (Fig. 11c), respectively. The residual stress of coating A is relatively large, and it is expressed as tensile stress. Correspondingly, it can be found that there are 14 cracks in the A coating in Fig. 12a. The residual stress of the b coating is half that of the A coating, which shows a slight tensile stress. In Fig. 12b, it can be seen that there are 6 cracks in the B coating surface. The residual stress of the C coating is further reduced than that of the B coating and also shows a slight tensile stress, corresponding to the 3 cracks that can be found in Fig. 12c. The method for evaluating the cracking rate of the cladding layer is generally expressed by the number of cracks per unit length or the crack length. The unit length used in the text is 100 mm. The cracking rates of A, B, and C coatings were 14, 6, and 3, respectively. The temperature field inhomogeneity during the laser cladding process and the large temperature gradient formed by the fluid flow in the molten pool, As a result, the temperature of the cladding layer after solidification is higher than that of the substrate. When the coating is cooled to room temperature, the shrinkage of the coating is larger than that of the substrate. The substrate is compressed and the cladding layer is drawn, which leads to the formation of cracks [24]. Fig. 13 shows the morphology of the fracture at the crack of the coating. It can be seen from Fig. 13a that there are many small tearing edges and secondary cracks on the fracture surface of coating A. at the same time, there are many small cleavage steps on the fracture surface, and there are fluvial patterns on the step surface. It is concluded that the fracture surface is brittle fracture and the cladding fracture is brittle fracture. Fig. 13b shows a small tear edge at the fracture of the B coating, with a small cleaving step. Fig. 13c shows the appearance of river patterns at the cracks of the C coating, cleavage steps and crystal planes, no secondary cracks were found on the crystal faces, and the possibility of brittle secondary phases was ruled out. In the process of cooling and solidification of laser cladding, the segregation of the alloying elements may form the residual phase of the low-melting
(2)
where W0 is standard sample wear loss, W1 is test material wear loss. It can be calculated that the relative wear resistance of A, B, and C
Fig. 9. Friction coefficients of the substrate and coatings versus time. 344
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(a)
(b)
200 ȝm
200 ȝm
(c)
(d)
200 ȝm
200 ȝm
Fig. 10. Micrographs of the worn surface (a) A coating (b) B coating (c) C coating (d) substrate.
current density increases, indicating that the oxide film on the surface of the coating is damaged due to anodic polarization, resulting in dissolution of the coating. The electrochemical parameters calculated from Tafel plots are listed in Table 1. The corrosion potential (Ecorr), corrosion current density (icorr), polarization resistance (Rp), and the anodic/ cathodic Tafel constants (Ba, Bc) were extracted directly from the potentiodynamic polarization curves by the Tafel fit method. As can be seen from the results of Table 1, from the corrosion speed point of view, v(Substrate) > v(A coating) > v(B coating) > v(C coating); from the corrosion kinetics point of view, icorr(C coating) < icorr(B coating) < icorr(A coating) < icorr(Substrate), Generally, Lower icorr, positive Ecorr, and higher Rp indicate that the coating provides better corrosion resistance. Therefore, the smaller the TiC particle size and the greater the concentration, the better the corrosion resistance of the coating. Fig. 15 is a Nyquist plot of the coating and substrate immersed for half an hour in a 3.5% NaCl solution. where Z’ is the real part of the impedance and Z” is the imaginary part. It can be seen that coating B and coating C show a higher total impedance, and the maximum impedance of the coatings can reach about 45,000 Ω. In this case, a large capacitance arc occurs in the high frequency region, and the radius of
eutectic. At the same time, the tensile stress due to shrinkage may cause intercrystalline microcracks in the unsolidified liquid phase. The crystal lattice that has grown up hinders the flow of liquid metal, which leads to the formation of hot cracks. The cracks caused by the residual stress were primarily crystal cracks, which belong tohotcracks, due to the liquid film formed by low melting point eutectic (residual liquid) weakened the bonding between the grains and producing crack under the action of residual stress. As a result, the greater the tensile stress was, the greater the produced cracks number were [25]. 3.8. Corrosion resistance analysis Fig. 14 shows the potentiodynamic polarization curves of the coatings and substrates in a 3.5% NaCl solution. The corrosion potentials of coating B and coating C are both close to −0.8 V. When the corrosion potential ranges from −0.8 to −0.4 V, the curve clearly indicates a passivation phenomenon. When the potential reaches −0.3 V, the passivation film on the coating surface is broken down, so that the corrosion current rapidly rises. The corrosion potentials of coating A are close to −0.9 V. When the polarization potential exceeds −0.9 V, the
Fig. 11. Surface Residual stress analysis of the coatings (a) A coating (b) B coating (c) C coating. 345
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Fig. 12. Crack distribution of the coating surface (a) A coating (b) B coating (c) C coating.
(a)
(c)
(b)
50 ȝm
50 ȝm
50 ȝm
Fig. 13. Fracture morphology of coating cracks (a) A coating (b) B coating (c) C coating.
Fig. 15. Nyquist plots of EIS of substrate and coatings.
Fig. 14. Potentiodynamic polarization of substrate and coatings.
substrate shows two capacitive arcs in the Nyquist plot, with a smaller radius at high frequencies and a larger radius at low frequencies [26]. According to the characteristics of electrochemical corrosion, the equivalent circuit shown on Fig. 16 is used to model the results from impedance testing the substrate and coatings. Phase elements are characterized by Q and n, and the CPE is considered to be the ideal capacitance when the value of n is 1. In the equivalent circuit of Fig. 16a, Rs is the resistance of the NaCl solution and, Rt is the resistance of the solution-substrate interface in parallel with CPEt. on Fig. 16b, Rs is the resistance of the NaCl solution, Rb is the transfer resistance of the coating in parallel with the constant phase element CPEb, and Rt is the barrier layer resistance of the coating in parallel with CPEt. the corresponding equivalent circuit parameters are listed in Table 2. The values
Table 1 Fitting data of polarization curves related to substrate and coatings. Sample
Ecorr(V)
Icorr(A·cm−2)
Ba(mv)
Bc(mv)
Rp(Ω·cm2)
A Coating B Coating C Coating Substrate
−0.92848 −0.80756 −0.75631 −1.01072
4.6941 × 10−7 5.5547 × 10−7 4.5395 × 10−8 2.7706 × 10−6
155.72 350.36 233.99 108.21
74.12 154.67 156.76 314.51
46,412 64,586 89,790 1294.7
the capacitor loop is large, so the corrosion resistance is better. Coating A exhibit capacitive arcs in the high-frequency region; the capacitive arc with a smaller radius corresponds to poor corrosion resistance. The 346
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Fig. 16. Equivalent circuits of the EIS plots for substrate (a) and coating (b). Table 2 Fitting data of EIS related to substrate and coatings. Sample
Rs (Ω.cm2)
Qb (Ω−1s−ncm−2)
Nb
Rb (kΩ.cm2)
Qt (Ω−1s−ncm−2)
Nt
Rt (kΩ.cm2)
Substrate A coating B coating C coating
4.057 4.185 4.651 12.91
– 5.48 × 10−4 3.432 × 10−4 2.417 × 10−6
– 0.8369 0.835 1
– 1.892 2.172 10.84
3.366 × 10−4 1.174 × 10−5 8.737 × 10−5 4.433 × 10−5
1 0.9311 1 0.8533
2.671 3.348 10.31 23.32
of Rt for coatings A, B and C were 3.348 kΩ·cm2, 10.31 kΩ·cm2, 23.32 kΩ·cm2, respectively, the coating is significantly better than the values of Rt for substrate. At the same time, the capacitance of the coating is obviously one order of magnitude lower than that of the substrate, which can improve the corrosion protection of the substrate [27,28]. Compared with the other coatings, the corrosion resistance of coating C is the best.
Jiangsu Higher Education Institutions (PAPD). References [1] H. Clemens, H. Kestler, Processing and applications of intermetallic γ-TiAl-based alloys, Adv. Eng. Mater. 2 (2000) 551–570. [2] F. Appel, U. Brossmann, U. Christoph, S. Eggert, P. Janschek, U. Lorenz, J. Müllauer, M. Oehring, J.D.H. Paul, Recent progress in the development of gamma titanium aluminide alloys, Adv. Eng. Mater. 2 (2000) 699–720. [3] J. Cheng, F. Li, S.Y. Zhu, Y. Yu, Z.H. Qiao, J. Yang, Electrochemical corrosion and tribological evaluation of TiAl alloy for marine application, Tribol. Int. 115 (2017) 483–549. [4] H.P. Qu, H.M. Wang, Microstructure and mechanical properties of laser melting deposited γ-TiAl intermetallic alloys, Mater. Sci. Eng. A 466 (2007) 187–194. [5] H. Wang, J. Cao, J. Feng, Brazing mechanism and infiltration strengthening of CC composites to TiAl alloys joint, Scripta Mater. 63 (2010) 859–862. [6] Y. Chen, H.M. Wang, Microstructure and high-temperature wear resistance of a laser surface alloyed γ-TiAl with carbon, Appl. Surf. Sci. 220 (2003) 186–192. [7] B. Cárcel, A. Serrano, J. Zambrano, V. Amigó, A.C. Cárcel, Laser cladding of TiAl intermetallic alloy on Ti6Al4V-process optimization and properties, Phys. Procedia. 56 (2014) 284–293. [8] R. Huang, J. Luo, C. Wang, Applications and developments of laser cladding technology, Ord. Mat. Sci. Eng. 4 (2005) 018. [9] F. Weng, C. Chen, H. Yu, Research status of laser cladding on titanium and its alloys: a review, Mater Design. 58 (2014) 412–425. [10] Y. Chen, H.M. Wang, Microstructure and wear resistance of a laser clad TiC reinforced nickel aluminides matrix composite coating, Mater. Sci. Eng. A 368 (2004) 80–87. [11] L. Chai, C. Wang, L. Han, L. Zhang, Influence of nano-TiC/C content on microstructure and properties of laser clad Ni-based alloy coatings, Chin. J. Lasers 41 (2014) 70–76. [12] 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. Tech. 200 (2006) 4462–4470. [13] J. Li, C. Chen, T. Squartini, Q. He, A study on wear resistance and microcrack of the Ti3Al/TiAl+TiC ceramic layer deposited by laser cladding on Ti–6Al–4V alloy, Appl. Surf. Sci. 257 (2010) 1550–1555. [14] A. Emamian, S.F. Corbin, A. Khajepour, The influence of combined laser parameters on in-situ formed TiC morphology during laser cladding, Surf. Coat. Tech. 206 (2011) 124–131. [15] Y.H. Liu, J. Li, F.Z. Xuan, Fabrication of TiC reinforced Ni based coating by laser cladding, Sur. Eng. 28 (2013) 560–563. [16] H. Wang, D. Gu, Nanometric TiC reinforced AlSi10Mg nanocomposites: powder preparation by high-energy ball milling and consolidation by selective laser melting, J. Compos. Mater. 49 (2015). [17] S. Yang, W. Liu, M. Zhong, Z. Wang, TiC reinforced composite coating produced by powder feeding laser cladding, Mater. Lett. 58 (2004) 2958–2962. [18] M. Parsons, E.V. Stepanov, A. Hiltner, E. Baer, Effect of strain rate on stepwise fatigue and creep slow crack growth in high density polyethylene, J. Mater. Sci. 35 (2000) 1857–1866. [19] A. Eltayeb, R.K. Vijayaraghavan, A.P. Mccoy, J. Cullen, S. Daniels, E. McGlynn, Control of crystal structure, morphology and optical properties of ceria films by post deposition annealing treatments, Thin Solid Films 603 (2016) 363–370. [20] R. Sun, Y. Lei, Microstructure and hardness of laser clad SiC β–Al composite
4. Conclusions (1) The morphology of the reinforced phase TiC in A coating is large, the distribution of TiC is scattered. A large number of micropores and cracks appear on the surface of the coating, and the cladding quality was poor. The morphology of B coating is connected granular and long structure, the distribution of TiC is uniform, the growth direction is regular, there were some fine micropores on the surface of the coating, and no cracks occurred. The morphology of the reinforced phase TiC in the C coating is granular and thin rod, and the granular TiC accumulates near the thin rod TiC and grows according to a certain rule. The dendrite arm in the cladding layer is finer, the microstructure growth is developed and the cladding quality is better. The dilution rates of A, B and C coatings are 9.09%, 6.54% and 5.21%, respectively. (2) The average microhardness of A, B, and C coatings were 598.04 HV0.2, 627.86HV0.2, and 701.51 HV0.2, respectively. The hardness of the coatings is increased by about 39%, 45% and 63% than that of the substrate, respectively. The wear mechanism of coating is mainly micro-cutting and scratching. The relative wear resistance of the A, B, and C coatings were 1.14, 1.17, 1.21, respectively, and C coating shows the best wear resistance. (3) The residual stresses of A, B, and C coatings were 143.3 MPa, 69.7 MPa, and 54.1 MPa, respectively. Both of them were tensile stress. The cracking rates of each coating were 14, 6, and 3, and the crack type was a crystalline crack in a hot crack. (4) When the particle size of TiC is small and the concentration is high, the corrosion resistance of the coating is relatively good. Acknowledgements The authors gratefully acknowledge the financial support from the Key Research and Development Project of Jiangsu Province (BE2016052) and the Priority Academic Program Development of 347
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coatings on Al alloys, Mater. Lett. 62 (2008) 3272–3275. [21] T. Chen, D. Liu, F. Wu, H. Wang, Effect of CeO2 on microstructure and wear resistance of TiC bioinert coatings on Ti6Al4V alloy by laser cladding, Materials 11 (2018) 58. [22] S. Yang, W. Liu, M. Zhong, Z. Wang, H. Kokawa, Fabrication of in-situ synthesized TiC particles reinforced composite coating by powder feeding laser cladding, J. Mater. Sci. 40 (2005) 2751–2754. [23] H. Zhou, P. Zhang, N. Sun, C.T. Wang, P.Y. Lin, L.Q. Ren, Wear properties of compact graphite cast iron with bionic units processed by deep laser cladding WC, Appl. Surf. Sci. 256 (2010) 6413–6419. [24] C. Lee, H. Park, J. Yoo, C. Lee, W.C. Woo, Residual stress and crack initiation in
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Full length article
Effects of TiC on the microstructure and properties of TiC/TiAl composite coating prepared by laser cladding
T
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X. Hea,b,c, R.G. Songa,b,c, , D.J. Kongb,d a
School of Materials Science and Engineering, Changzhou University, Changzhou, Jiangsu 213164, China Jiangsu Key Laboratory of Materials Surface Science and Technology, Changzhou University, Changzhou, Jiangsu 213164, China c Jiangsu Collaborative Innovation Center of Photovolatic Science and Engineering, Changzhou University, Changzhou, Jiangsu 213164, China d School of Mechanical Engineering, Changzhou University, Changzhou, Jiangsu 213164, China b
H I GH L IG H T S
prepared TiC/TiAl composite coatings on TiAl alloy by laser cladding. • We microstructure of coating containing nano-TiC was finer and less defects. • The coating can improve the hardness and wear resistance of the substrate. • TiC/TiAl • The coating can significantly improve the corrosion resistance of the substrate.
A R T I C LE I N FO
A B S T R A C T
Keywords: TiAl alloy Laser cladding TiC/ TiAl composite coating Microstructure Properties
In order to study the effect of TiC on the microstructure and properties of cladding coatings, TiC/ TiAl composite coatings have been prepared by using a laser cladding technique on TiAl alloy, and the microstructure and properties of the resulting composite coatings have been investigated using scanning electron microscopy (SEM), X-ray diffraction (XRD), micro-hardness testing, X-ray stress measurements, friction and wear testing, and an electrochemical workstation. The results showed that when the particle size of TiC is up to micron, the growth of TiC is more developed, the dendrite growth direction is disordered, and the cladding quality is poor. When the particle size of TiC is nanoscale, the morphology of the reinforcement phase TiC is mainly granular and thin rods, which is uniform in the coating and the direction of growth is regular. When the concentration of TiC reaches 20%, the microstructure of the cladding layer grows well and is dense. The microhardness and wear resistance of the coatings prepared by adding nano TiC are better than those of the coatings prepared by micron TiC. When the content of nano TiC in the coatings increases from 10% to 20%, the microhardness and wear resistance of the coatings are obviously improved. The residual stress of the coating is positively related to the crack rate. The residual stress in the coating is tensile stress, and the crack type is the crystal crack in the hot crack. When the particle size of TiC is small and the concentration is high, the coating shows better corrosion resistance.
1. Introduction TiAl alloy is a new kind of intermetallic compound structure material, which has many outstanding characteristics, such as low density, high specific strength and specific modulus, high specific rigidity and good high-temperature creep resistance. The combination of these advantages will push TiAl alloy to be used in automotive, aerospace and marine industries. Therefore, TiAl intermetallic compounds have attracted wide attention. However, TiAl alloys display poor high-temperature oxidation resistance as high-temperature parts. Moreover, the
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low hardness and poor tribological properties severely limit their applications as key moving components. When exposed to the sea water environment, they are not immune to combined corrosion and wear attacks. Although protective passive films with a thickness of 3–5 μm readily form on the surface of titanium alloy in sea water, they are too thin to resist the shear stress and removed quickly, accelerating the metal dissolution and corrosion. In order to overcome this drawback, the formations of ceramic coatings on the TiAl surface are proposed and proven to be the effective methods to improve the surface properties of alloys. Among various materials, the TiAl intermetallic based coating is
Corresponding author at: School of Materials Science and Engineering, Changzhou University, Changzhou, Jiangsu 213164, China. E-mail address: [email protected] (R.G. Song).
https://doi.org/10.1016/j.optlastec.2018.11.037 Received 4 July 2018; Received in revised form 21 October 2018; Accepted 18 November 2018 0030-3992/ © 2018 Elsevier Ltd. All rights reserved.
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residual stress were as follows: fixed peak method for cross correlation, using a Co target, Co-Kα1 radiation source, Prague crystal face of (4 0 0), incident angles of 0°, 25°, 35°, and 45°, stress constant of −130 MPa/°, 2θ scan start and termination angles of 155–145°, 2θ scanning step of 0.10°, counting time of 0.50 s, X-ray light tube high voltage of 22 kV, and an X-ray light tube current of 6.0 mA. Electrochemical tests were carried out by using CS350 type electrochemical workstation. The sample dimension of electrochemical test was 10 mm × 10 mm × 3 mm. Except for the surface of the coating, the rests of the surfaces were inlaid with epoxy resin. The test medium was a 3.5% NaCl solution. The system of electrochemical workstation included a saturate calomel electrode (SCE) as a reference electrode, Pt as a counter electrode, sample as working electrodes. The potentiodynamic scanning rate of 1 mV/s, sampling frequency of 0.5 Hz, the test potential range of −0.5 to 0.5 V, and test time of 1800 s. EIS measurement's frequency range is 10−1 to 105, and test time of 300 s. All the measurements were performed after 30 min immersion in 3.5% NaCl solution until open circuit potential (OCP) was stable.
a promising method to enhance the high-temperature oxidation resistance, corrosion and tribological properties of TiAl alloys [1–3]. In recent years, in order to further study TiAl alloys, many researchers have improved the wear resistance, high-temperature oxidation resistance and corrosion resistance of TiAl intermetamerics by a series of surface modification techniques such as ion infiltration, deposition, laser alloying, and laser cladding [4–7]. Laser cladding is an advanced surface modification technology. The thickness of the cladding layer obtained by laser cladding is large, and the structure is dense and uniform [8–9]. Reinforced carbides TiC possess both excellent hightemperature abrasive and corrosion resistance because of their very high hardness, unique strong atomic bonds and high stability under high-temperature exposure. So the high-temperature coatings reinforced by TiC carbides are expected to possess good high-temperature corrosive wear resistance [10]. At present, Chai et al. studied the effect of nano-TiC/C on the microstructure and properties of laser cladding Ni-base alloy coating. The results show that with the increase of nanoTiC/C addition, the number of γ-Ni dendrites showed a trend of increasing first and then decreasing, and the change trend of its size was reversed [11]. Liu has also studied the preparation of γ/W2C/TiC composite coatings on TiAl alloys. The results show that the wear resistance and high-temperature oxidation resistance of the coatings are significantly improved [12]. At present, there are few studies at home and abroad to study the effect of particle size and concentration of TiC on the microstructure and properties of laser cladding composite coatings [13–15]. Therefore, this paper adopts laser cladding technology to prepare TiC-reinforced composite coatings on TiAl alloys, and analyzes the influence of TiC on the structure, hardness, wear resistance and corrosion resistance of cladding layers.
3. Results and discussion 3.1. XRD and morphologies analysis of powders Fig. 1 shows the X-ray diffraction pattern of the mixed powder used for the laser cladding test. Because of the lower milling speed, the mixed powder did not produce a new phase during ball milling, and it was still Ti, Al and TiC phases [16]. Fig. 2 shows the micromorphology of TiC powder at nanometer and micrometer diameters. The nano TiC used in the experiment is spherical and has a particle size of about 50 nm. The micro TiC used is in the form of a block or a sheet with irregular morphology. The particle size is about 40 μm, which contains fine TiC debris.
2. Experimental The experimental material and contrast material was TiAl alloys, which are mainly composed of Ti-48Al-2Cr-2Nb (at.%). The sample size is φ46 × 9 mm. Before the experiment, the specimens were sanded step by step with sandpaper and washed repeatedly with anhydrous ethanol. The spherical powders used in the cladding test are mainly Ti powder (purity 99.3%, 64 μm–120 μm) and Al powder (purity 99.0%, 220 μm–330 μm). Ti powder and Al powder are mixed according to the atomic ratio of 1:1, then three different TiC powders were added. 20% micron TiC was added to the group A powder, 10% nano-TiC were added to the group B powder, and 20% nano-TiC were added to the group C powder. The coating was chemically pre-prepared, and methylcellulose was used as a binder. The thickness was about 1 mm, and it was then incubated in a drying oven at 80 °C. for 2 h. The laser power was fixed at 2.40 KW, the scanning speed was 10 mm/s, the multi-pass rate is 40%, the laser cladding adopts a circular spot, the diameter is 3 mm, the above process parameters are fixed, and the preparation contains 20% TiC (μm) A coating, 10% TiC (nm) B coating, 20% TiC (nm) C coating. The morphologies, element compositions and phases of coatings were analyzed using a JSM-6360LA type scanning electron microscope (SEM), energy dispersive spectrometer (EDS), and D/max2500 PC X-ray diffraction (XRD), respectively. The microhardness of the coating was measured with an THL200 digital microhardness tester with the following conditions: a loading of 200 g and a loading time of 15 s. The hardness of cladding layer is measured every 100 μm from the cladding surface to the substrate, and the average value is obtained three times at the same depth level. Friction and wear testing were conducted using a CFT-1 surface tester. The grinding material used was SiC at a loading of 150 N and a motor speed of 400 r/min, using a reciprocating sliding mode, the wear scar radius was 3 mm and the running time was 30 min. The wear medium was air and the measurement was completed using a BT25S electronic analytical balance to measure the weight loss. The residual stresses of the coatings were analyzed using an X-ray stress tester. The technological parameters used for measurement of the
3.2. Morphologies of coating interfaces Fig. 3 shows the metallographic cross-section of the coating. It can be seen that the cross-section of the cladding layer is in turn divided into four zones: the coating zone (CZ), the diffusion zone (DZ), the heat affected zone (HAZ) and the substrate. In the A coating, the microstructure of the cladding layer is dense, and the thickness of the cladding layer is about 0.8 mm. The reinforcing phase TiC is scattered in the cladding layer and has no obvious gradient distribution characteristics [17]. In the B coating, the structure of the cladding layer is sparse, the thickness is about 0.75 mm, and the reinforcing phase TiC is more regularly distributed in the coating, and it can be seen that the
Fig. 1. XRD patterns of the mixed powder. 340
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(b)
(a)
100 nm
100 ȝm
(c)
(d)
Fig. 2. Morphology of TiC with different particle sizes and the diameter distribution of the TiC powder. (a, c) nano TiC (b, d) micron TiC.
considered that the dilution rate of the coating is better at 5%.
transition between the cladding layer, and the substrate is natural. The microstructure of C coating was well-developed and dense. The thickness of cladding layer was about 0.7 mm. The microstructure of the coating grows up regularly, and the TiC distribution of the reinforcing phase is more uniform. Because the dilution rate directly affects the performance of the coating, the dilution rate λ is expressed by the formula
λ=
h × 100% H+h
3.3. XRD analysis of coating surfaces Fig. 4 shows the XRD pattern of the surface of the coatings obtained by using A, B and C three groups of powders at P = 2.40 kW and V = 10 mm/s. It can be seen that the major components of the A coating are For TiAl and TiC phases, compared with B and C coatings, Ti3Al does not appear at about 40° for A coating, and the main composition phase for B and C coatings is TiAl and TiC phases and a small amount of Ti3Al phase.
(1)
where H is the coating thickness and h is the melting depth of substrate. The melting depth of A, B and C coatings is 0.082 mm, 0.052 mm, 0.039 mm respectively. So the dilution rates of coatings A, B and C can be calculated as 9.03%, 6.51% and 5.18%, respectively. It is generally
Fig. 3. Micrographs of coating cross-section (a) A coating (b) B coating (c) C coating. 341
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d, the main elements in the cladding layer are Ti, Al and C. The atomic ratio of Ti to Al is about 2:1. Therefore, in addition to the TiAl phase, the cladding layer should also contain Ti3Al phases. This is consistent with the results of XRD pattern analysis in Fig. 4. Therefore, the main composition phase in the coating is the continuous phase TiAl. TiC is the reinforcing phase and also contains a small amount of Ti3Al phase. Due to the small TiC content in B powder, the growth phase of TiC grows slowly, and at the same time. The nano TiC in B powder has a spherical shape and regular morphology, resulting in a more uniform distribution in the cladding layer and a more regular growth direction. There are some tiny pores on the surface of the coating, no cracks appear, and the cladding quality is good. Fig. 5c shows the typical microstructure of C coating. Combined with the results of energy spectrum analysis of Fig. 6e and f, the main elements in the cladding layer are Ti, Al, C, and the atomic ratio of Ti to Al is about 2:1, The continuous phase is TiAl phase and a small amount of Ti3Al phase in the C coting. The reinforcing phase TiC is mainly in the form of particles, fine rods. Granular TiC gathers next to the thin rod-shaped TiC, grows according to a certain rule, the dendrite arms in the cladding layer are finer. The organization grows well. Compared with the B coating, the TiC mass fraction in the C coating is 10% greater than that of the B coating, resulting in more dense TiC in the C coating. The surface pores of the coating are further reduced, so that the quality of the cladding layer is improved. Combined with the above test analysis results, it can be found that the coating quality of the A coating is poor, and the coating quality of the B and C coatings is better. Because the microstructure of the crystal has a decisive influence on its growth morphology and mode [19], when the TiC particle size in the A coating and the B and C coatings is greatly different, the A coating is exposed to a high-energy laser beam. Micron TiC absorbs less energy than nano-TiC in the other two coatings, causing TiC in A coating to not fully melt, affecting the temperature field and element distribution in the molten pool, making TiC provide a nucleation surface for crystals in the molten pool, further causing heat stress concentration, the formation of cracks and holes. Comparing B and C coatings, the microstructure and growth direction are not much different. Granular TiC and thin rod-shaped TiC are connected with each other according to certain rules. However, the reinforcement phase TiC is denser and has fewer micropores on the surface of the C coating, which is undoubtedly due to TiC with more C coatings.
Fig. 4. XRD patterns of cladding layer.
3.4. Morphologies and EDS analysis of coating surfaces Figs. 5 and 6 shows the typical microstructures and their energy spectrum analysis results for A, B and C coatings, respectively. It can be seen from Fig. 5a that the reinforcement phase TiC growth of the A coating is relatively well-developed, the structure of the coating is larger and the direction of dendrite growth is disorderly. It is not typical. At the same time, it can be seen that there are a large number of tiny pores and a small number of cracks on the surface of the cladding layer, the coating quality is poor. On the one hand, it may be due to the irregular shape of TiC in the powder A and the small fines that cause stress concentration and cracks on the top of the pores. On the other hand, the thermal physics coefficient between micro TiC and the matrix is very different. This can easily lead to the generation of thermal stress, resulting in cracks and pores, making the cladding less effective [18]. According to the energy spectrum analysis of Fig. 6a and 6b, the main elements in the coating are Ti, Al and C. The ratio of Ti and Al atoms are close to 1:1. Therefore, the continuous phase (Spectrum 1) in the cladding layer is mainly TiAl. Which is mainly due to irregular shape of TiC in A powder, poor fluidity, large particle size, and more laser energy required for cladding, thus affecting its temperature field and element field distribution in the molten pool, making they do not diffuse to each other at the bottom of the bath and do not reach the temperature needed to form the liquid phase, and thus the Ti3Al phase cannot be formed. The reinforced phase in the cladding layer is TiC (Spectrum 2), which is consistent with the XRD pattern analysis results in Fig. 4. Fig. 5b shows a typical tissue morphology of the B coating. The dendrites in the cladding layer are mostly connected in a granular or elongated form. According to energy spectrum analysis of Fig. 6c and
3.5. Microhardness analysis The microhardness distribution from the cladding surface to the substrate is shown on Fig. 7. The microhardness of C coating increases slowly from 0 to 500 μm, and the microhardness reaches a maximum of 813.3 HV0.2 at a distance of about 500 μm from the surface, then the microhardness gradually decreases during the extension to the
(b)
(a)
(c)
Spectrum1
Spectrum 5
Spectrum 3
Crack Spectrum 2
Pore
50 ȝm
Spectrum 6
Spectrum 4
50 ȝm
50 ȝm
Fig. 5. The surface SEM images of the TiC/TiAl coating (a) A coating (b) B coating (c) C coating. 342
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Fig. 6. EDS spectrum of cladding layer of the tested TiC/TiAl coating (a, b) A coating (c, d) B coating (e, f) C coating.
microhardness of the substrate is about 430 HV0.2, and the average microhardness of the A, B, and C coatings are 701.51 HV0.2, 627.86 HV0.2, and 598.04 HV0.2, respectively. Compared with the substrate, the coating has a significantly increased microhardness, especially for the C coating, the microhardness is increased by nearly 50% due to the rich reinforcing phase TiC in the coating surface. However, it can be clearly observed that the microhardness of the B and C coatings is better than that of the A coating. On the other hand, since the B and C coatings are nano reinforcing phase TiC, the resulting microstructure is finer and the fine-grain strengthening effect is more pronounced. In addition, B and C coatings contain Ti3Al phase, which has close-packed hexagonal superlattice structure, high specific strength and high elastic modulus. which can significantly improve the surface hardness of the cladding layer [20].
3.6. Friction and wear testing Fig.7. Microhardness distribution of the TiC/TiAl coatings.
Fig. 8 shows the data of the wear loss tests for the three composite coatings and substrates, it can be seen that the wear of the A, B and C coatings are decreased in varying degrees compared to the original TiAl alloy, the wear loss of the C coating is significantly better than the A and B coatings. On the one hand, due to the distribution of a large number of nanoparticle reinforced phases TiC in the C coating, it has a
substrate. The microhardness of B coating gradually increases from 0 to 300 μm, and reaches a maximum of 798.3HV0.2 at a distance of about 300 μm. The microhardness of A coating reaches a maximum of 669.1 HV0.2 at 300 μm, and then gradually decreases. The average 343
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coatings are 1.14, 1.17, and 1.21, respectively. Therefore, the hardness and wear resistance of the cladding layer can be greatly improved when the TiC concentration of is higher and the grain size is smaller [22]. Fig. 10 demonstrates the surface morphologies of the coatings and substrate afte wearing test. Fig. 10a shows the wear profile of the A coating. It can be seen that the worn surface is rough and distributed with more grooves. The surface has a layer of shallow furrows, and the wear mechanism is mainly micro-cutting and scratching. Fig. 10b shows the wear profile of the B coating. It can be observed that the surface of the coating is relatively flat, with some hard particles on the surface. There are some tiny pits and scratches, but it is not obvious. This is mainly due to the hardness of the B coating. Higher, the cutting effect on the coating weakens. On the other hand, the hard reinforcing phase TiC protrudes on the surface of the substrate and reduces the contact surface with the grinding wheel, but as the wear progresses, the protruding TiC will fall off from the coating surface, and thus Leave pits. The wear surface of the C coating is smoother than that of the B coating, with only minor scratches and partial damage. The wear surface is basically free of plastic deformation and adhesive wear. As the high hardness of nano-TiC diffuses on the surface of the coating, microcutting during grinding is further weakened, and therefore only slight scratches are generated. At the same time. The presence of high hardness TiC makes it play a role in the hard skeleton support of the cladding layer, thereby effectively reducing the wear rate [23]. The wear surface of the substrate is very uneven as shown in Fig. 10d. After the wear, the furrow is deep and wide, there are many traces of friction, the direction is disordered, and there are shallow peels locally. The wear mechanism is mainly abrasive wear.
Fig. 8. Total wear loss of the TiC/TiAl coatings.
higher hardness and strength, and it plays a role in abrasion against in the process of rubbing with the friction pair and hinders the plastic deformation of the coated substrate, and helps to improve the wear resistance of the coating. On the other hand, when the molten pool solidifies at high speed, some elements such as Cr and Nb were infiltrated into the substrate, which played a solid solution strengthening role in the coating. Meanwhile, the dense structure and lower cracks and pores in the B and C coatings are also important factors to enhance the wear resistance of the coating [21]. With the passage of time, the wear loss appears to be gradually increasing. This is due to the high brittleness of TiAl alloys at room temperature. The shedding of TiC inclusions in the wear process forms a new abrasive between the two pairs of grinding parts, and TiC has high hardness, which can be cut into the cladding layer and make the wear increase. Fig. 9 shows the air friction coefficient as a function of time for coatings and substrate, it can be seen from Fig. 9 that the friction coefficient increases rapidly to about 0.6 in 0–3 min, and finally the friction coefficient stabilizes at about 0.5. Compared with the A coating, due to the higher hardness of the B and C coatings, the friction coefficient is further reduced, and finally stabilized at 0.45 and 0.35, respectively. The relative wear resistance is one of the important criteria for judging frictional wear performance, and the relative wear resistance (ε) is defined as
ε=
W0 W1
3.7. Residual stress analysis The cracks were related with residual stresses in the Laser cladding test, from the 2θ-sin2φ regression curves, the residual stress of Cladding coating A, B and C was 143.3 MPa (Fig. 11a), 69.7 MPa (Fig. 11b), 54.1 MPa (Fig. 11c), respectively. The residual stress of coating A is relatively large, and it is expressed as tensile stress. Correspondingly, it can be found that there are 14 cracks in the A coating in Fig. 12a. The residual stress of the b coating is half that of the A coating, which shows a slight tensile stress. In Fig. 12b, it can be seen that there are 6 cracks in the B coating surface. The residual stress of the C coating is further reduced than that of the B coating and also shows a slight tensile stress, corresponding to the 3 cracks that can be found in Fig. 12c. The method for evaluating the cracking rate of the cladding layer is generally expressed by the number of cracks per unit length or the crack length. The unit length used in the text is 100 mm. The cracking rates of A, B, and C coatings were 14, 6, and 3, respectively. The temperature field inhomogeneity during the laser cladding process and the large temperature gradient formed by the fluid flow in the molten pool, As a result, the temperature of the cladding layer after solidification is higher than that of the substrate. When the coating is cooled to room temperature, the shrinkage of the coating is larger than that of the substrate. The substrate is compressed and the cladding layer is drawn, which leads to the formation of cracks [24]. Fig. 13 shows the morphology of the fracture at the crack of the coating. It can be seen from Fig. 13a that there are many small tearing edges and secondary cracks on the fracture surface of coating A. at the same time, there are many small cleavage steps on the fracture surface, and there are fluvial patterns on the step surface. It is concluded that the fracture surface is brittle fracture and the cladding fracture is brittle fracture. Fig. 13b shows a small tear edge at the fracture of the B coating, with a small cleaving step. Fig. 13c shows the appearance of river patterns at the cracks of the C coating, cleavage steps and crystal planes, no secondary cracks were found on the crystal faces, and the possibility of brittle secondary phases was ruled out. In the process of cooling and solidification of laser cladding, the segregation of the alloying elements may form the residual phase of the low-melting
(2)
where W0 is standard sample wear loss, W1 is test material wear loss. It can be calculated that the relative wear resistance of A, B, and C
Fig. 9. Friction coefficients of the substrate and coatings versus time. 344
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(a)
(b)
200 ȝm
200 ȝm
(c)
(d)
200 ȝm
200 ȝm
Fig. 10. Micrographs of the worn surface (a) A coating (b) B coating (c) C coating (d) substrate.
current density increases, indicating that the oxide film on the surface of the coating is damaged due to anodic polarization, resulting in dissolution of the coating. The electrochemical parameters calculated from Tafel plots are listed in Table 1. The corrosion potential (Ecorr), corrosion current density (icorr), polarization resistance (Rp), and the anodic/ cathodic Tafel constants (Ba, Bc) were extracted directly from the potentiodynamic polarization curves by the Tafel fit method. As can be seen from the results of Table 1, from the corrosion speed point of view, v(Substrate) > v(A coating) > v(B coating) > v(C coating); from the corrosion kinetics point of view, icorr(C coating) < icorr(B coating) < icorr(A coating) < icorr(Substrate), Generally, Lower icorr, positive Ecorr, and higher Rp indicate that the coating provides better corrosion resistance. Therefore, the smaller the TiC particle size and the greater the concentration, the better the corrosion resistance of the coating. Fig. 15 is a Nyquist plot of the coating and substrate immersed for half an hour in a 3.5% NaCl solution. where Z’ is the real part of the impedance and Z” is the imaginary part. It can be seen that coating B and coating C show a higher total impedance, and the maximum impedance of the coatings can reach about 45,000 Ω. In this case, a large capacitance arc occurs in the high frequency region, and the radius of
eutectic. At the same time, the tensile stress due to shrinkage may cause intercrystalline microcracks in the unsolidified liquid phase. The crystal lattice that has grown up hinders the flow of liquid metal, which leads to the formation of hot cracks. The cracks caused by the residual stress were primarily crystal cracks, which belong tohotcracks, due to the liquid film formed by low melting point eutectic (residual liquid) weakened the bonding between the grains and producing crack under the action of residual stress. As a result, the greater the tensile stress was, the greater the produced cracks number were [25]. 3.8. Corrosion resistance analysis Fig. 14 shows the potentiodynamic polarization curves of the coatings and substrates in a 3.5% NaCl solution. The corrosion potentials of coating B and coating C are both close to −0.8 V. When the corrosion potential ranges from −0.8 to −0.4 V, the curve clearly indicates a passivation phenomenon. When the potential reaches −0.3 V, the passivation film on the coating surface is broken down, so that the corrosion current rapidly rises. The corrosion potentials of coating A are close to −0.9 V. When the polarization potential exceeds −0.9 V, the
Fig. 11. Surface Residual stress analysis of the coatings (a) A coating (b) B coating (c) C coating. 345
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Fig. 12. Crack distribution of the coating surface (a) A coating (b) B coating (c) C coating.
(a)
(c)
(b)
50 ȝm
50 ȝm
50 ȝm
Fig. 13. Fracture morphology of coating cracks (a) A coating (b) B coating (c) C coating.
Fig. 15. Nyquist plots of EIS of substrate and coatings.
Fig. 14. Potentiodynamic polarization of substrate and coatings.
substrate shows two capacitive arcs in the Nyquist plot, with a smaller radius at high frequencies and a larger radius at low frequencies [26]. According to the characteristics of electrochemical corrosion, the equivalent circuit shown on Fig. 16 is used to model the results from impedance testing the substrate and coatings. Phase elements are characterized by Q and n, and the CPE is considered to be the ideal capacitance when the value of n is 1. In the equivalent circuit of Fig. 16a, Rs is the resistance of the NaCl solution and, Rt is the resistance of the solution-substrate interface in parallel with CPEt. on Fig. 16b, Rs is the resistance of the NaCl solution, Rb is the transfer resistance of the coating in parallel with the constant phase element CPEb, and Rt is the barrier layer resistance of the coating in parallel with CPEt. the corresponding equivalent circuit parameters are listed in Table 2. The values
Table 1 Fitting data of polarization curves related to substrate and coatings. Sample
Ecorr(V)
Icorr(A·cm−2)
Ba(mv)
Bc(mv)
Rp(Ω·cm2)
A Coating B Coating C Coating Substrate
−0.92848 −0.80756 −0.75631 −1.01072
4.6941 × 10−7 5.5547 × 10−7 4.5395 × 10−8 2.7706 × 10−6
155.72 350.36 233.99 108.21
74.12 154.67 156.76 314.51
46,412 64,586 89,790 1294.7
the capacitor loop is large, so the corrosion resistance is better. Coating A exhibit capacitive arcs in the high-frequency region; the capacitive arc with a smaller radius corresponds to poor corrosion resistance. The 346
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Fig. 16. Equivalent circuits of the EIS plots for substrate (a) and coating (b). Table 2 Fitting data of EIS related to substrate and coatings. Sample
Rs (Ω.cm2)
Qb (Ω−1s−ncm−2)
Nb
Rb (kΩ.cm2)
Qt (Ω−1s−ncm−2)
Nt
Rt (kΩ.cm2)
Substrate A coating B coating C coating
4.057 4.185 4.651 12.91
– 5.48 × 10−4 3.432 × 10−4 2.417 × 10−6
– 0.8369 0.835 1
– 1.892 2.172 10.84
3.366 × 10−4 1.174 × 10−5 8.737 × 10−5 4.433 × 10−5
1 0.9311 1 0.8533
2.671 3.348 10.31 23.32
of Rt for coatings A, B and C were 3.348 kΩ·cm2, 10.31 kΩ·cm2, 23.32 kΩ·cm2, respectively, the coating is significantly better than the values of Rt for substrate. At the same time, the capacitance of the coating is obviously one order of magnitude lower than that of the substrate, which can improve the corrosion protection of the substrate [27,28]. Compared with the other coatings, the corrosion resistance of coating C is the best.
Jiangsu Higher Education Institutions (PAPD). References [1] H. Clemens, H. Kestler, Processing and applications of intermetallic γ-TiAl-based alloys, Adv. Eng. Mater. 2 (2000) 551–570. [2] F. Appel, U. Brossmann, U. Christoph, S. Eggert, P. Janschek, U. Lorenz, J. Müllauer, M. Oehring, J.D.H. Paul, Recent progress in the development of gamma titanium aluminide alloys, Adv. Eng. Mater. 2 (2000) 699–720. [3] J. Cheng, F. Li, S.Y. Zhu, Y. Yu, Z.H. Qiao, J. Yang, Electrochemical corrosion and tribological evaluation of TiAl alloy for marine application, Tribol. Int. 115 (2017) 483–549. [4] H.P. Qu, H.M. Wang, Microstructure and mechanical properties of laser melting deposited γ-TiAl intermetallic alloys, Mater. Sci. Eng. A 466 (2007) 187–194. [5] H. Wang, J. Cao, J. Feng, Brazing mechanism and infiltration strengthening of CC composites to TiAl alloys joint, Scripta Mater. 63 (2010) 859–862. [6] Y. Chen, H.M. Wang, Microstructure and high-temperature wear resistance of a laser surface alloyed γ-TiAl with carbon, Appl. Surf. Sci. 220 (2003) 186–192. [7] B. Cárcel, A. Serrano, J. Zambrano, V. Amigó, A.C. Cárcel, Laser cladding of TiAl intermetallic alloy on Ti6Al4V-process optimization and properties, Phys. Procedia. 56 (2014) 284–293. [8] R. Huang, J. Luo, C. Wang, Applications and developments of laser cladding technology, Ord. Mat. Sci. Eng. 4 (2005) 018. [9] F. Weng, C. Chen, H. Yu, Research status of laser cladding on titanium and its alloys: a review, Mater Design. 58 (2014) 412–425. [10] Y. Chen, H.M. Wang, Microstructure and wear resistance of a laser clad TiC reinforced nickel aluminides matrix composite coating, Mater. Sci. Eng. A 368 (2004) 80–87. [11] L. Chai, C. Wang, L. Han, L. Zhang, Influence of nano-TiC/C content on microstructure and properties of laser clad Ni-based alloy coatings, Chin. J. Lasers 41 (2014) 70–76. [12] 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. Tech. 200 (2006) 4462–4470. [13] J. Li, C. Chen, T. Squartini, Q. He, A study on wear resistance and microcrack of the Ti3Al/TiAl+TiC ceramic layer deposited by laser cladding on Ti–6Al–4V alloy, Appl. Surf. Sci. 257 (2010) 1550–1555. [14] A. Emamian, S.F. Corbin, A. Khajepour, The influence of combined laser parameters on in-situ formed TiC morphology during laser cladding, Surf. Coat. Tech. 206 (2011) 124–131. [15] Y.H. Liu, J. Li, F.Z. Xuan, Fabrication of TiC reinforced Ni based coating by laser cladding, Sur. Eng. 28 (2013) 560–563. [16] H. Wang, D. Gu, Nanometric TiC reinforced AlSi10Mg nanocomposites: powder preparation by high-energy ball milling and consolidation by selective laser melting, J. Compos. Mater. 49 (2015). [17] S. Yang, W. Liu, M. Zhong, Z. Wang, TiC reinforced composite coating produced by powder feeding laser cladding, Mater. Lett. 58 (2004) 2958–2962. [18] M. Parsons, E.V. Stepanov, A. Hiltner, E. Baer, Effect of strain rate on stepwise fatigue and creep slow crack growth in high density polyethylene, J. Mater. Sci. 35 (2000) 1857–1866. [19] A. Eltayeb, R.K. Vijayaraghavan, A.P. Mccoy, J. Cullen, S. Daniels, E. McGlynn, Control of crystal structure, morphology and optical properties of ceria films by post deposition annealing treatments, Thin Solid Films 603 (2016) 363–370. [20] R. Sun, Y. Lei, Microstructure and hardness of laser clad SiC β–Al composite
4. Conclusions (1) The morphology of the reinforced phase TiC in A coating is large, the distribution of TiC is scattered. A large number of micropores and cracks appear on the surface of the coating, and the cladding quality was poor. The morphology of B coating is connected granular and long structure, the distribution of TiC is uniform, the growth direction is regular, there were some fine micropores on the surface of the coating, and no cracks occurred. The morphology of the reinforced phase TiC in the C coating is granular and thin rod, and the granular TiC accumulates near the thin rod TiC and grows according to a certain rule. The dendrite arm in the cladding layer is finer, the microstructure growth is developed and the cladding quality is better. The dilution rates of A, B and C coatings are 9.09%, 6.54% and 5.21%, respectively. (2) The average microhardness of A, B, and C coatings were 598.04 HV0.2, 627.86HV0.2, and 701.51 HV0.2, respectively. The hardness of the coatings is increased by about 39%, 45% and 63% than that of the substrate, respectively. The wear mechanism of coating is mainly micro-cutting and scratching. The relative wear resistance of the A, B, and C coatings were 1.14, 1.17, 1.21, respectively, and C coating shows the best wear resistance. (3) The residual stresses of A, B, and C coatings were 143.3 MPa, 69.7 MPa, and 54.1 MPa, respectively. Both of them were tensile stress. The cracking rates of each coating were 14, 6, and 3, and the crack type was a crystalline crack in a hot crack. (4) When the particle size of TiC is small and the concentration is high, the corrosion resistance of the coating is relatively good. Acknowledgements The authors gratefully acknowledge the financial support from the Key Research and Development Project of Jiangsu Province (BE2016052) and the Priority Academic Program Development of 347
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