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Microstructures, mechanical properties and corrosion resistance of Hastelloy C22 coating produced by laser cladding
Journal of Alloys and Compounds 553 (2013) 253–258
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Microstructures, mechanical properties and corrosion resistance of Hastelloy C22 coating produced by laser cladding Qin-Ying Wang a, Yang-Fei Zhang a, Shu-Lin Bai a,⇑, Zong-De Liu b a b
Department of Materials Science and Engineering, HEDPS, Center for Applied Physics and Technology, LTCS, College of Engineering, Peking University, Beijing 100871, China Key Laboratory of Condition Monitoring and Control for Power Plant Equipment of Ministry of Education, North China Electric Power University, Beijing 102206, China
a r t i c l e
i n f o
Article history: Received 19 August 2012 Accepted 25 October 2012 Available online 16 November 2012 Keywords: Laser cladding Hastelloy C22 coating Microstructure Mechanical property Corrosion resistance
a b s t r a c t The Hastelloy C22 coatings H1 and H2 were prepared by laser cladding technique with laser scanning speeds of 6 and 12 mm/s, respectively. Their microstructures, mechanical properties and corrosion resistance were investigated. The microstructures and phase compositions were studied by metallurgical microscope, scanning electron microscope and X-ray diffraction analysis. The hardness and scratch resistance were measured by micro-hardness and nanoindentation tests. The polarization curves and electrochemical impedance spectroscopy were tested by electrochemical workstation. Planar, cellular and dendritic solidifications were observed in the coating cross-sections. The coatings metallurgically wellbonded with the substrate are mainly composed of primary phase c-nickel with solution of Fe, W, Cr and grain boundary precipitate of Mo6Ni6C. The hardness and corrosion resistance of steel substrate are significantly improved by laser cladding Hastelloy C22 coating. Coating H2 shows higher microhardness than that of H1 by 34% and it also exhibits better corrosion resistance. The results indicate that the increase of laser scanning speed improves the microstuctures, mechanical properties and corrosion resistance of Hastelloy C22 coating. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Steel is easily corroded in damp environments of seawater, chemical plant, soil and underground. Surface coating deposited on the steel has become an effective way to improve the substrate properties, such as corrosion resistance, wear and oxidation behaviors. Some deposition techniques have been developed to produce the coatings with excellent mechanical and anticorrosion properties, including vacuum arc melting, plasma spraying and laser cladding [1,2]. Laser cladding is widely chosen to produce metallurgically well-bonded coating by the advantages of high power, small heat-affected-zone (HAZ) and minimal distortion of the substrate [3–6]. The traditional laser sources are Nd: YAG and CO2. In recent years, diode laser cladding technique with higher power and lower cost has been developed to create beam profile. Kennedy et al. [7] reported that diode laser exhibited better modal stability than other laser sources. Zhu et al. [8] found that the porosity of AZ31 welds by diode laser welding is less than that by CO2 laser welding. The technological parameters of laser cladding are important to the coating quality, including laser power, spot diameter, overlapping ratio, etc. For example, dramatic fluctuation of temperature during laser cladding might cause porosities
⇑ Corresponding author. Tel.: +86 10 6275 9379; fax: +86 10 6275 1812. E-mail address: [email protected] (S.-L. Bai). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.10.193
and cracks [9] and high laser scanning speed would aggravate elemental dilution between coating and substrate [10]. The influence of laser cladding parameters on the microstructure and properties is different for various types of materials. Hamedi et al. [11] found that the micro-hardness of TiC coating was increased as pulse duration increased from 4 to 12 ms. Viejo et al. [12] reported that the optimum laser scanning speed for A380/SiC/10p composite with highest corrosion potential was 80 mm/s. Volovitch et al. [13] found that minimizing the microstructure inhomogeneity by using higher laser scanning speed could improve the corrosion resistance of Al–Si coating. Hastelloy C22 is a nickel-based alloy with proper concentrations of Cr, Mo and W [14,15]. It is one of the few materials that are resistant to the corrosion of hydrogen chloride, hypochlorite and chlorine dioxide, which can be applied for protection and decoration of various constructions, marine facilities, machines and, etc. [16,17]. However, the high cost of Hastelloy C22 limits its large-scale applications. In recent decade, Hastelloy C22 coating has been rapidly developed. Haemers et al. [18] produced Hastelloy coating with a thickness of 3 mm by laser cladding technique in 2001. Vignolo et al. [19] reported a synthesis of Hastelloy thin film by XeCl pulsed laser ablation in 2002. Then Zocco et al. [1] found that the corrosion behavior of Hastelloy coating by XeCl pulsed laser ablation approximated to the bulk one. However, there is still lack of study reported on Hastelloy C22 coating produced by high power diode continuous laser cladding technique.
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The present work aims at developing Hastelloy C22 coating on widely used Q235 steel substrate with preset powders method and by high power diode continuous laser cladding technique. Two laser scanning speeds of 6 and 12 mm/s were chosen to avoid the large deformation of steel substrate under high power. The mechanical properties and corrosion resistance of coatings were investigated by divers methods and discussed according to the features of microstructures. 2. Materials and experimental methods 2.1. Materials studied and laser cladding process The chemical compositions of Hastelloy C22 powders (self-preparation) and Q235 steel substrate are given in Table 1. By particle size analysis, the powder size range is 30–50 lm. A layer of well-mixed Hastelloy C22 powders with a thickness of 1 mm was preset on the substrate before laser cladding. A continuous diode laser system (RuiChi-4000, China) was used to carry out the cladding with laser power of 3 kW, spot diameter of 8 mm and overlap ratio of 35%. Two scanning speeds of 6 and 12 mm/s were adopted, corresponding to the coatings named of H1 and H2, respectively. Argon was used as the laser shielding gas to produce oxide free coatings. Laser beam was perpendicular to the substrate surface. After laser cladding, the as-prepared coating/substrate bulk was cut into a size of 11.5 11.5 4 mm3. The surface and cross-section of samples were polished with 2000# abrasive papers and corundum powders of 3–5 lm, then ultrasonically cleaned in distilled water followed by ethanol and acetone. The thickness of each coating after surface treatment was 0.8 mm. The samples for electrochemical tests were covered with silicagel except an exposure coating area of 10 10 mm2. 2.2. Experimental methods The microstructures and phase compositions of the Hastelloy C22 coatings were measured by using scanning electron microscope (SEM, S-4800 HITACHI), metallurgical microscope (Olympus, BX51M) and X-ray diffractometer (XRD, DMAX-2400). The micro-hardness of coating cross-section was measured by a Vickers hardness tester (HMV-2T, Japan) under a load of 0.98 N and dwell time of 15 s. Triboindenter system (Hysitron Inc., USA) with a Berkovich tip of a nominal radius of curvature equal to 50 nm was used to measure nano-scale hardness of the coating top surface. The maximum penetration load was 2000 lN with the loading and unloading rate of 200 lN/s. The holding segment of 2 s at the maximum load was undertaken to eliminate the effect of creep deformation. Twenty indents were carried out on each coating to obtain the accurate average hardness. A 60° cono spherical diamond tip of 1 lm was used in nano-scratch test with a constant load of 1000 lN and velocity of 0.33 lm/s. During the experiments, the temperature was controlled at 20 °C and relative humidity at 32%. Electrochemical measurements were carried out on an electrochemical workstation (CHI760, China) at 25 °C with the samples immersed in 3.5 wt.% NaCl solution. Each test was repeated at least three times to obtain reliable result. The potentiodynamic polarization curves were measured at the scanning speed of 1 mV/s, commencing from 1 V to 0 V. The electrochemical impedance spectroscopy (EIS) was measured during 16 days under the frequency range from 100 KHz to 10 mHz, and voltage amplitude of ±5 mV rms. The conventional threeelectrode cell was used, with an Ag/AgCl/saturated KCl electrode as the reference electrode, a platinum electrode as the counter electrode and the coating sample as the working electrode.
3. Results and discussion 3.1. Microstructure and composition The microstructures of coatings H1 and H2 cross-sections (not near the ovelap region) with few defects of porosity and crack are shown in Fig. 1. The coating/substrate interface of both coatings is regular and well-bonded, and the structure of each coating is mainly composed of eutectic and dentrite. Three types of solidifications, planar (the bottom connecting the substrate), cellular
(the middle) and dendritic (the upside near the top surface) successively occurred at the cross sections. This morphology mainly depends on the ratio G/R from interface to coating surface, where R is the solidification rate and G is the gradient of temperature [20]. At the interface, the ratio G/R is very large, and there is hardly any ingredient super cooling, so the planar solidification is dominant. As the ratio G/R decreases, the ingredient super cooling turns up and develops gradually, which causes cellular solidification to occur nearby the planar region, and then dendritic solidification after cellular one in both coatings. From the interface to coating center, cellular and dendritic growth is perpendicular to the interface due to heat transmittance by the substrate. The microstructures become very fine with decreasing interdendritic spacing near the top surface for the high solidification rate. Besides, compared with coating H1, the size of columnar dendrite in coating H2 is reduced and so its microstructure is compact due to the decline of the primary dendrite arm spacing caused by higher laser scanning speed of 12 mm/s. This can be expressed by the equation below [21]:
pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi pffiffiffiffiffiffi k ¼ a gP =½ðT T 0 Þ 2kV s cos h
ð1Þ
where k is the primary dendrite arm spacing, a is coefficient, g is laser absorption coefficient, P is laser power, T is liquid temperature of the alloy, T0 is the initial temperature of substrate, k is the thermal conductivity of material, Vs is laser scanning speed and h is the angle between laser scanning direction and solidification direction. Therefore, k would decline as Vs increases, which leads to compact and relatively small grains in coating H2. On the contrary, the coarsening columnar dendrite becomes the microstructural characteristic of coating H1 with laser scanning speed of 6 mm/s. The dendritic solidification in the overlap region of coating H1 top surface develops at an angle of 60° to the laser cladding direction, as shown in Fig. 2. This is due to the remelting and resolidification in the overlap region caused by the adjacent laser cladding process. The remelting makes the previously formed cellular solidification become the pre-nucleus for the columnar dendritic growth in resolidification. The growth direction of columnar dendritic solidification is determined by the heat flux direction. However, this directional growth is not observed in H2. Fig. 3 shows schematically the mechanisms of microstructure difference in the overlap regions of coatings H1 and H2. Theoretically, the heat of the melt pool could be mostly conducted to the substrate along C1(Z) direction due to the great temperature difference between coating and substrate for the single-pass cladding. In this case, the solidification develops perpendicular to the coating/substrate interface and opposite to the direction of the heat dissipation. However, for the multi-track cladding in this study, the heat of the overlap region remelted by adjacent laser cladding could be conducted not only to the substrate C1(Z), but also to previously deposited region along C2(X) and the trailing end of the current laser melt pool along C3(Y). Then synthesizing the three possibilities of heat dissipation, the final solidification is opposite to the general heat dissipation direction (C00 ), and at an angle h to the laser cladding direction on the top surface (XY plane). This angle h should be larger than 45° but smaller than 90° for the reason of more heat dissipation along C2 (fully cooling direction) than C3 (relatively slow cooling direction). Dinda et al. [22] found the similar phenomenon in the interface region between different
Table 1 Chemical compositions of Hastelloy C22 powders and Q235 steel.
Powders Q235 steel
Elements
Cr
Mo
Fe
W
Ni
Content (wt.%) Content (wt.%)
21.00
13.00
5.00 Bal.
5.00
Bal.
Si
Mn
S
P
C
0.37
0.08
0.04
0.04
0.16
Q.-Y. Wang et al. / Journal of Alloys and Compounds 553 (2013) 253–258
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Fig. 1. Cross-section SEM micrographs of coatings H1 and H2.
Fig. 2. Metallurgical microscope micrographs of overlap regions in coatings H1 and H2 top surfaces.
layers of the metal coating produced by laser cladding technique. On the contrary, few directional dendritic growth but fine and small equiaxed crystals are observed in coating H2 overlap region. In our opinion, it is related with high cooling rate, leaving not enough time for previous cellular solidification to develop into the dendritic ones. XRD patterns of coatings H1 and H2 are shown in Fig. 4. It can be seen that both coatings contain c-nickel solid solution as the primary phase. The diffraction peaks of c-nickel solid solution moved a little towards small angle campared with the XRD standard diffraction card due to the solution of elements Cr, W, Fe, etc.. However, relatively distinct peaks of the hard phase of
Fig. 3. Schematic presentation of the mechanism of directional dendritic growth.
Mo6Ni6C was detected in the diffraction pattern of coating H2. This is due to the high cooling rate of melt pool when the diode laser beams move as fast as 12 mm/s, leading to the formation of metastable phase. Corresponding to Fig. 1, the dentrite consists of the primary phase c-nickel solid solution, which is surrounded by the eutectic mainly composed of Mo6Ni6C phase [23]. 3.2. Mechanical properties 3.2.1. Vickers hardness The micro-hardness variation of coatings H1 and H2 at the cross section from coating to substrate is shown in Fig. 5. The abrupt change of micro-hardness at the coating/substrate interface indicates the reinforcement of coating on the hardness of the substrate. The micro-hardness of coatings H1 and H2 is about 40% and 60% higher than that of substrate, respectively. Besides, the average micro-hardness of coating H2 is 34% higher than that of coating H1. This is due to, in one hand, the existence of large quantity of Mo6Ni6C hard phase in coating H2 and its compact microstructure in other hand. 3.2.2. Nano-scale hardness The nano-scale hardness of constitutive microstructures in Hastelloy C22 coatings can be measured by nanoindentation technique [24]. The load–displacement curves of structures eutectic and dentrite in both coatings are shown in Fig. 6. The average hardness of eutectic in coatings H1 and H2 is 10.71 (H1-E) and 11.73 GPa (H2-E), and that of dentrite is 6.56 (H1-D) and 6.94 GPa (H2-D), respectively. The higher hardness of eutectic
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Fig. 4. XRD patterns of coatings H1 and H2.
Fig. 6. Load–displacement curves of nanoindentation on coatings H1 and H2.
compared to dentrite is due to the combined strengthening effect of precipitates and hard phase. In fact, the average nano-scale hardness of eutectic or dentrite in both coatings is not much different, yet this is not the case of Vickers hardness. The reason is that the small size indenter tip of nanoindentation acts only on a single structure, and the response of the structure to the indentation reflect its hardness. While by micro-scale Vickers tests, the domain on which indenter acts is so large that a group of dentrite and eutectic are pressed at the same time, so the hard eutectic significantly influences Vickers hardness. Actually the size-effect plays the key role in the hardness measurements, i.e. nanoindentation gives the single structure hardness, while Vickers hardness is that of bulk material. The volume fractions of eutectic calculated by particle size analysis of coatings H1 and H2 are 7.72% and 12.16%, respectively. The average nano-scale hardness of both coatings H1 and H2 are predicted using the simple rule of mixture with the following formula [25]:
ðHn Þav ¼ ðHn Þe V f þ ðHn Þd ð1 V f Þ
ð2Þ
where (Hn)av is the average hardness of coating, (Hn)e and (Hn)d are nano-scale hardness of eutectic and dentrite measured by nanoindentation tests, Vf is the volume fraction of eutectic in the coating. The average hardness of coatings H1 and H2 is shown in Table 2. Coating H2 exhibits higher hardness of 7.52 GPa than coating H1 of 6.88 GPa. This is consistent with the Vickers hardness of both coatings. Although the values of hardness measured by Vickers
Table 2 Hardness of coatings H1 and H2 by nanoindentation tests and prediction.
Coating H1 Coating H2
(Hn)e/GPa
(Hn)d/GPa
Ratio of eutectic/%
(Hn)av/GPa
10.71 11.73
6.56 6.94
7.72 12.16
6.88 7.52
hardness tester and nanoindentation system could not be compared with each other for the different definition and operating mode, the variation tendency of hardness will be valid. 3.2.3. Nano-scratch tests Nano-scratch images at the substrate, coatings H1 and H2 top surfaces by nanoindentation system are presented in Fig. 7. Three phenomena are observed in terms of the groove morphology. First, the groove of substrate is deeper and wider than that of coatings H1 and H2, which means that the coatings are more resistant to nano-scratch than substrate; second, with more concentration of Mo6Ni6C hard phase in the eutectic of coating H2, its scratch groove is more shallow with hardly any ridge volume, showing better scratch resistance than that of coating H1; third, the eutectic and dentrite in each coating display notably different ridge volumes, showing clear concurrence with the nano-scale hardness measured by nanoindentition. The deeper scratch groove at dentrite with more ridge volume than that at eutentic in each coating indicates relatively low scratch resistance of dentrite. The frictional behaviors of substrate, coatings H1 and H2 are obtained from nano-scratch tests. The coefficient of friction indicates the resistance of material to tip penetration in the normal direction, which is determined by the ratio of the lateral and normal force [26,27]. The similar curves of friction coefficient and lateral force with sliding distance are shown in Fig. 8. This indicates that the friction coefficient is determined by lateral force due to the constant normal force of 1000 lN used in this work. The coating H2 has smallest general friction coefficient around 0.26, while they are 0.33 and 0.39 for coating H1 and substrate, indicating best tribological property of coating H2. Besides, the lower friction coefficient of eutectic than that of dentrite is found in each coating, i.e. the decline of friction coefficient curve when the tip scratched across the eutectic. 3.3. Electrochemical measurements
Fig. 5. Vickers hardness along the distance from coating to substrate of coatings H1 and H2.
3.3.1. Polarization curves The polarization curves of substrate and coatings H1, H2 are plotted in Fig. 9. The results indicate that the corrosion resistance
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Fig. 7. Images of the nano-scratch grooves at substrate, coatings H1 and H2.
Fig. 9. Polarization curves of substrate, coatings H1 and H2.
Fig. 8. Variation of lateral force and friction coefficient with sliding distance.
of both coatings is much better than that of substrate. The most important reason is the different chemical compositions, i.e. more than 98% Fe in steel substrate makes it easily corroded. Instead, for the Hastelloy C22 coatings, the dense Cr2O3 layer is formed on the coating surface, preventing the coating from corrosion. Moreover, coating H2 displays better corrosion resistance with higher corrosion potential (Ecorr) of 0.286 V and lower corrosion current density (Icorr) of 2.45 107 A/cm2, than coating H1 with the values of 0.301 V and 3.40 107 A/cm2. Coating H2 has more compact microstructure and smaller grain size than coating H1 as stated in Fig. 1, thus there are more grain boundary and dislocations in coating H2, which provide more active nucleation sites for quick formation of continuous and protective passive film. The passive film plays a key role in preventing Ni ions or electrons from migrating towards the surface to enhance electrochemical reaction and aggravate corrosion. There was very little weight loss of coatings H1 and H2 with shining surface after immersing in 3.5 wt.% NaCl solution for several months. The fact indicates that the anticorrosion property of Hastelloy C22 coating prepared by laser cladding is as good as that of bulk one. The weight loss rate related to corrosion current density (Icorr) can be roughly predicted through the assumption of uniform corrosion and Faraday’s law [28]:
Q ¼ DW n N A =M
ð3Þ
DW ¼ W=At ¼ iM=nF
ð4Þ
Table 3 Corrosion results of the substrate, Hastelloy C22 coatings H1 and H2. Samples
H1 coating
H2 coating
Substrate
Average atomic mass (g/mol) Icorr (A/cm2) DW (mg/cm2 s) P (mm/y)
58.33 3.40 107 6.85 108 0.0025
60.43 2.45 10-7 5.11 108 0.0019
56 1.45 105 4.31 106 0.1792
where Q = total charges (Coulomb), DW = weight loss rate (g/cm2 s), M = atomic mass of the metal (g/mol), n = number of electron/atoms consumed (around 3 for both coatings, 2 for substrate), NA = Avogadro constant (6.022 1023), A = exposure area (10 10 mm2 in this paper), t = corrosion time (s), i = current density (A/cm2), F = Faraday’s constant (96,485 Coulomb/mol). Furthermore, through the relation below, the change of weight loss rate can also be converted to the average penetration rate:
P ¼ 3:65DW=q
ð5Þ
where P is the average penetration rate (mm/y), q is coating density (8.69 g/cm3 for both coatings). The final corrosion results are given in Table 3, which illustrates that coating H2 demonstrates higher corrosion resistance than coating H1 and substrate. 3.3.2. EIS measurements Electrochemical impedance spectroscopy (EIS) was measured to evaluate the variation of electrochemical reactions in the substrate and coatings. The Nyquist plots of the substrate during 7 days and
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Fig. 10. EIS of substrate, coatings H1 and H2.
coatings H1 and H2 during 16 days in 3.5 wt.% NaCl solution are shown in Fig. 10. The substrate exhibits one capacitive semicircle in the whole frequency range, and then the semicircle turns small and depresses with the increase of immersing time. This reveals that the local corrosion might have occurred on the substrate. Comparatively, the Nyquist plots of coatings show the continuously increasing a quarter capacitive circle with immersing time and then tend to stabilize, indicating the gradual formation of stable passive film. Besides, larger capacitive quarter-circle radius for coating H2 than for coating H1 is found, which reveals more compact and thicker passive film in coating H2. 4. Conclusions The Hastelloy C22 coatings H1 and H2 with few defects were produced by laser cladding with scanning speeds of 6 and 12 mm/s, respectively. Planar, cellular and dendritic solidifications were observed on the cross-sections of both coatings. Coating H2 shows more compact microstructure than that of coating H1 due to the higher laser scanning speed. Higher micro-hardness was found in coating H2 than H1 due to the existence of large quantity of eutectic containing Mo6Ni6C hard phase and the compact microstructure. For the same reason, coating H2 exhibits better tribological property. The corrosion resistance of coating H2, characterized with high corrosion potential of 0.286 V and low corrosion current density of 2.45 107 A/cm2, is better than H1 with the values of 0.301 V and 3.40 107 A/cm2 due to the quick formation of thick passive film. The results reveal that the microstructures, mechanical properties and corrosion resistance of Hastelloy C22 coatings are influenced by laser scanning speed. References [1] A. Zocco, A. Perrone, M.F. Vignolo, S. Duhalde, I. Avram, C. Morales, T. Pérez, Appl. Surf. Sci. 208–209 (2003) 669–675.
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Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Microstructures, mechanical properties and corrosion resistance of Hastelloy C22 coating produced by laser cladding Qin-Ying Wang a, Yang-Fei Zhang a, Shu-Lin Bai a,⇑, Zong-De Liu b a b
Department of Materials Science and Engineering, HEDPS, Center for Applied Physics and Technology, LTCS, College of Engineering, Peking University, Beijing 100871, China Key Laboratory of Condition Monitoring and Control for Power Plant Equipment of Ministry of Education, North China Electric Power University, Beijing 102206, China
a r t i c l e
i n f o
Article history: Received 19 August 2012 Accepted 25 October 2012 Available online 16 November 2012 Keywords: Laser cladding Hastelloy C22 coating Microstructure Mechanical property Corrosion resistance
a b s t r a c t The Hastelloy C22 coatings H1 and H2 were prepared by laser cladding technique with laser scanning speeds of 6 and 12 mm/s, respectively. Their microstructures, mechanical properties and corrosion resistance were investigated. The microstructures and phase compositions were studied by metallurgical microscope, scanning electron microscope and X-ray diffraction analysis. The hardness and scratch resistance were measured by micro-hardness and nanoindentation tests. The polarization curves and electrochemical impedance spectroscopy were tested by electrochemical workstation. Planar, cellular and dendritic solidifications were observed in the coating cross-sections. The coatings metallurgically wellbonded with the substrate are mainly composed of primary phase c-nickel with solution of Fe, W, Cr and grain boundary precipitate of Mo6Ni6C. The hardness and corrosion resistance of steel substrate are significantly improved by laser cladding Hastelloy C22 coating. Coating H2 shows higher microhardness than that of H1 by 34% and it also exhibits better corrosion resistance. The results indicate that the increase of laser scanning speed improves the microstuctures, mechanical properties and corrosion resistance of Hastelloy C22 coating. Ó 2012 Elsevier B.V. All rights reserved.
1. Introduction Steel is easily corroded in damp environments of seawater, chemical plant, soil and underground. Surface coating deposited on the steel has become an effective way to improve the substrate properties, such as corrosion resistance, wear and oxidation behaviors. Some deposition techniques have been developed to produce the coatings with excellent mechanical and anticorrosion properties, including vacuum arc melting, plasma spraying and laser cladding [1,2]. Laser cladding is widely chosen to produce metallurgically well-bonded coating by the advantages of high power, small heat-affected-zone (HAZ) and minimal distortion of the substrate [3–6]. The traditional laser sources are Nd: YAG and CO2. In recent years, diode laser cladding technique with higher power and lower cost has been developed to create beam profile. Kennedy et al. [7] reported that diode laser exhibited better modal stability than other laser sources. Zhu et al. [8] found that the porosity of AZ31 welds by diode laser welding is less than that by CO2 laser welding. The technological parameters of laser cladding are important to the coating quality, including laser power, spot diameter, overlapping ratio, etc. For example, dramatic fluctuation of temperature during laser cladding might cause porosities
⇑ Corresponding author. Tel.: +86 10 6275 9379; fax: +86 10 6275 1812. E-mail address: [email protected] (S.-L. Bai). 0925-8388/$ - see front matter Ó 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2012.10.193
and cracks [9] and high laser scanning speed would aggravate elemental dilution between coating and substrate [10]. The influence of laser cladding parameters on the microstructure and properties is different for various types of materials. Hamedi et al. [11] found that the micro-hardness of TiC coating was increased as pulse duration increased from 4 to 12 ms. Viejo et al. [12] reported that the optimum laser scanning speed for A380/SiC/10p composite with highest corrosion potential was 80 mm/s. Volovitch et al. [13] found that minimizing the microstructure inhomogeneity by using higher laser scanning speed could improve the corrosion resistance of Al–Si coating. Hastelloy C22 is a nickel-based alloy with proper concentrations of Cr, Mo and W [14,15]. It is one of the few materials that are resistant to the corrosion of hydrogen chloride, hypochlorite and chlorine dioxide, which can be applied for protection and decoration of various constructions, marine facilities, machines and, etc. [16,17]. However, the high cost of Hastelloy C22 limits its large-scale applications. In recent decade, Hastelloy C22 coating has been rapidly developed. Haemers et al. [18] produced Hastelloy coating with a thickness of 3 mm by laser cladding technique in 2001. Vignolo et al. [19] reported a synthesis of Hastelloy thin film by XeCl pulsed laser ablation in 2002. Then Zocco et al. [1] found that the corrosion behavior of Hastelloy coating by XeCl pulsed laser ablation approximated to the bulk one. However, there is still lack of study reported on Hastelloy C22 coating produced by high power diode continuous laser cladding technique.
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The present work aims at developing Hastelloy C22 coating on widely used Q235 steel substrate with preset powders method and by high power diode continuous laser cladding technique. Two laser scanning speeds of 6 and 12 mm/s were chosen to avoid the large deformation of steel substrate under high power. The mechanical properties and corrosion resistance of coatings were investigated by divers methods and discussed according to the features of microstructures. 2. Materials and experimental methods 2.1. Materials studied and laser cladding process The chemical compositions of Hastelloy C22 powders (self-preparation) and Q235 steel substrate are given in Table 1. By particle size analysis, the powder size range is 30–50 lm. A layer of well-mixed Hastelloy C22 powders with a thickness of 1 mm was preset on the substrate before laser cladding. A continuous diode laser system (RuiChi-4000, China) was used to carry out the cladding with laser power of 3 kW, spot diameter of 8 mm and overlap ratio of 35%. Two scanning speeds of 6 and 12 mm/s were adopted, corresponding to the coatings named of H1 and H2, respectively. Argon was used as the laser shielding gas to produce oxide free coatings. Laser beam was perpendicular to the substrate surface. After laser cladding, the as-prepared coating/substrate bulk was cut into a size of 11.5 11.5 4 mm3. The surface and cross-section of samples were polished with 2000# abrasive papers and corundum powders of 3–5 lm, then ultrasonically cleaned in distilled water followed by ethanol and acetone. The thickness of each coating after surface treatment was 0.8 mm. The samples for electrochemical tests were covered with silicagel except an exposure coating area of 10 10 mm2. 2.2. Experimental methods The microstructures and phase compositions of the Hastelloy C22 coatings were measured by using scanning electron microscope (SEM, S-4800 HITACHI), metallurgical microscope (Olympus, BX51M) and X-ray diffractometer (XRD, DMAX-2400). The micro-hardness of coating cross-section was measured by a Vickers hardness tester (HMV-2T, Japan) under a load of 0.98 N and dwell time of 15 s. Triboindenter system (Hysitron Inc., USA) with a Berkovich tip of a nominal radius of curvature equal to 50 nm was used to measure nano-scale hardness of the coating top surface. The maximum penetration load was 2000 lN with the loading and unloading rate of 200 lN/s. The holding segment of 2 s at the maximum load was undertaken to eliminate the effect of creep deformation. Twenty indents were carried out on each coating to obtain the accurate average hardness. A 60° cono spherical diamond tip of 1 lm was used in nano-scratch test with a constant load of 1000 lN and velocity of 0.33 lm/s. During the experiments, the temperature was controlled at 20 °C and relative humidity at 32%. Electrochemical measurements were carried out on an electrochemical workstation (CHI760, China) at 25 °C with the samples immersed in 3.5 wt.% NaCl solution. Each test was repeated at least three times to obtain reliable result. The potentiodynamic polarization curves were measured at the scanning speed of 1 mV/s, commencing from 1 V to 0 V. The electrochemical impedance spectroscopy (EIS) was measured during 16 days under the frequency range from 100 KHz to 10 mHz, and voltage amplitude of ±5 mV rms. The conventional threeelectrode cell was used, with an Ag/AgCl/saturated KCl electrode as the reference electrode, a platinum electrode as the counter electrode and the coating sample as the working electrode.
3. Results and discussion 3.1. Microstructure and composition The microstructures of coatings H1 and H2 cross-sections (not near the ovelap region) with few defects of porosity and crack are shown in Fig. 1. The coating/substrate interface of both coatings is regular and well-bonded, and the structure of each coating is mainly composed of eutectic and dentrite. Three types of solidifications, planar (the bottom connecting the substrate), cellular
(the middle) and dendritic (the upside near the top surface) successively occurred at the cross sections. This morphology mainly depends on the ratio G/R from interface to coating surface, where R is the solidification rate and G is the gradient of temperature [20]. At the interface, the ratio G/R is very large, and there is hardly any ingredient super cooling, so the planar solidification is dominant. As the ratio G/R decreases, the ingredient super cooling turns up and develops gradually, which causes cellular solidification to occur nearby the planar region, and then dendritic solidification after cellular one in both coatings. From the interface to coating center, cellular and dendritic growth is perpendicular to the interface due to heat transmittance by the substrate. The microstructures become very fine with decreasing interdendritic spacing near the top surface for the high solidification rate. Besides, compared with coating H1, the size of columnar dendrite in coating H2 is reduced and so its microstructure is compact due to the decline of the primary dendrite arm spacing caused by higher laser scanning speed of 12 mm/s. This can be expressed by the equation below [21]:
pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi pffiffiffiffiffiffi k ¼ a gP =½ðT T 0 Þ 2kV s cos h
ð1Þ
where k is the primary dendrite arm spacing, a is coefficient, g is laser absorption coefficient, P is laser power, T is liquid temperature of the alloy, T0 is the initial temperature of substrate, k is the thermal conductivity of material, Vs is laser scanning speed and h is the angle between laser scanning direction and solidification direction. Therefore, k would decline as Vs increases, which leads to compact and relatively small grains in coating H2. On the contrary, the coarsening columnar dendrite becomes the microstructural characteristic of coating H1 with laser scanning speed of 6 mm/s. The dendritic solidification in the overlap region of coating H1 top surface develops at an angle of 60° to the laser cladding direction, as shown in Fig. 2. This is due to the remelting and resolidification in the overlap region caused by the adjacent laser cladding process. The remelting makes the previously formed cellular solidification become the pre-nucleus for the columnar dendritic growth in resolidification. The growth direction of columnar dendritic solidification is determined by the heat flux direction. However, this directional growth is not observed in H2. Fig. 3 shows schematically the mechanisms of microstructure difference in the overlap regions of coatings H1 and H2. Theoretically, the heat of the melt pool could be mostly conducted to the substrate along C1(Z) direction due to the great temperature difference between coating and substrate for the single-pass cladding. In this case, the solidification develops perpendicular to the coating/substrate interface and opposite to the direction of the heat dissipation. However, for the multi-track cladding in this study, the heat of the overlap region remelted by adjacent laser cladding could be conducted not only to the substrate C1(Z), but also to previously deposited region along C2(X) and the trailing end of the current laser melt pool along C3(Y). Then synthesizing the three possibilities of heat dissipation, the final solidification is opposite to the general heat dissipation direction (C00 ), and at an angle h to the laser cladding direction on the top surface (XY plane). This angle h should be larger than 45° but smaller than 90° for the reason of more heat dissipation along C2 (fully cooling direction) than C3 (relatively slow cooling direction). Dinda et al. [22] found the similar phenomenon in the interface region between different
Table 1 Chemical compositions of Hastelloy C22 powders and Q235 steel.
Powders Q235 steel
Elements
Cr
Mo
Fe
W
Ni
Content (wt.%) Content (wt.%)
21.00
13.00
5.00 Bal.
5.00
Bal.
Si
Mn
S
P
C
0.37
0.08
0.04
0.04
0.16
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Fig. 1. Cross-section SEM micrographs of coatings H1 and H2.
Fig. 2. Metallurgical microscope micrographs of overlap regions in coatings H1 and H2 top surfaces.
layers of the metal coating produced by laser cladding technique. On the contrary, few directional dendritic growth but fine and small equiaxed crystals are observed in coating H2 overlap region. In our opinion, it is related with high cooling rate, leaving not enough time for previous cellular solidification to develop into the dendritic ones. XRD patterns of coatings H1 and H2 are shown in Fig. 4. It can be seen that both coatings contain c-nickel solid solution as the primary phase. The diffraction peaks of c-nickel solid solution moved a little towards small angle campared with the XRD standard diffraction card due to the solution of elements Cr, W, Fe, etc.. However, relatively distinct peaks of the hard phase of
Fig. 3. Schematic presentation of the mechanism of directional dendritic growth.
Mo6Ni6C was detected in the diffraction pattern of coating H2. This is due to the high cooling rate of melt pool when the diode laser beams move as fast as 12 mm/s, leading to the formation of metastable phase. Corresponding to Fig. 1, the dentrite consists of the primary phase c-nickel solid solution, which is surrounded by the eutectic mainly composed of Mo6Ni6C phase [23]. 3.2. Mechanical properties 3.2.1. Vickers hardness The micro-hardness variation of coatings H1 and H2 at the cross section from coating to substrate is shown in Fig. 5. The abrupt change of micro-hardness at the coating/substrate interface indicates the reinforcement of coating on the hardness of the substrate. The micro-hardness of coatings H1 and H2 is about 40% and 60% higher than that of substrate, respectively. Besides, the average micro-hardness of coating H2 is 34% higher than that of coating H1. This is due to, in one hand, the existence of large quantity of Mo6Ni6C hard phase in coating H2 and its compact microstructure in other hand. 3.2.2. Nano-scale hardness The nano-scale hardness of constitutive microstructures in Hastelloy C22 coatings can be measured by nanoindentation technique [24]. The load–displacement curves of structures eutectic and dentrite in both coatings are shown in Fig. 6. The average hardness of eutectic in coatings H1 and H2 is 10.71 (H1-E) and 11.73 GPa (H2-E), and that of dentrite is 6.56 (H1-D) and 6.94 GPa (H2-D), respectively. The higher hardness of eutectic
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Fig. 4. XRD patterns of coatings H1 and H2.
Fig. 6. Load–displacement curves of nanoindentation on coatings H1 and H2.
compared to dentrite is due to the combined strengthening effect of precipitates and hard phase. In fact, the average nano-scale hardness of eutectic or dentrite in both coatings is not much different, yet this is not the case of Vickers hardness. The reason is that the small size indenter tip of nanoindentation acts only on a single structure, and the response of the structure to the indentation reflect its hardness. While by micro-scale Vickers tests, the domain on which indenter acts is so large that a group of dentrite and eutectic are pressed at the same time, so the hard eutectic significantly influences Vickers hardness. Actually the size-effect plays the key role in the hardness measurements, i.e. nanoindentation gives the single structure hardness, while Vickers hardness is that of bulk material. The volume fractions of eutectic calculated by particle size analysis of coatings H1 and H2 are 7.72% and 12.16%, respectively. The average nano-scale hardness of both coatings H1 and H2 are predicted using the simple rule of mixture with the following formula [25]:
ðHn Þav ¼ ðHn Þe V f þ ðHn Þd ð1 V f Þ
ð2Þ
where (Hn)av is the average hardness of coating, (Hn)e and (Hn)d are nano-scale hardness of eutectic and dentrite measured by nanoindentation tests, Vf is the volume fraction of eutectic in the coating. The average hardness of coatings H1 and H2 is shown in Table 2. Coating H2 exhibits higher hardness of 7.52 GPa than coating H1 of 6.88 GPa. This is consistent with the Vickers hardness of both coatings. Although the values of hardness measured by Vickers
Table 2 Hardness of coatings H1 and H2 by nanoindentation tests and prediction.
Coating H1 Coating H2
(Hn)e/GPa
(Hn)d/GPa
Ratio of eutectic/%
(Hn)av/GPa
10.71 11.73
6.56 6.94
7.72 12.16
6.88 7.52
hardness tester and nanoindentation system could not be compared with each other for the different definition and operating mode, the variation tendency of hardness will be valid. 3.2.3. Nano-scratch tests Nano-scratch images at the substrate, coatings H1 and H2 top surfaces by nanoindentation system are presented in Fig. 7. Three phenomena are observed in terms of the groove morphology. First, the groove of substrate is deeper and wider than that of coatings H1 and H2, which means that the coatings are more resistant to nano-scratch than substrate; second, with more concentration of Mo6Ni6C hard phase in the eutectic of coating H2, its scratch groove is more shallow with hardly any ridge volume, showing better scratch resistance than that of coating H1; third, the eutectic and dentrite in each coating display notably different ridge volumes, showing clear concurrence with the nano-scale hardness measured by nanoindentition. The deeper scratch groove at dentrite with more ridge volume than that at eutentic in each coating indicates relatively low scratch resistance of dentrite. The frictional behaviors of substrate, coatings H1 and H2 are obtained from nano-scratch tests. The coefficient of friction indicates the resistance of material to tip penetration in the normal direction, which is determined by the ratio of the lateral and normal force [26,27]. The similar curves of friction coefficient and lateral force with sliding distance are shown in Fig. 8. This indicates that the friction coefficient is determined by lateral force due to the constant normal force of 1000 lN used in this work. The coating H2 has smallest general friction coefficient around 0.26, while they are 0.33 and 0.39 for coating H1 and substrate, indicating best tribological property of coating H2. Besides, the lower friction coefficient of eutectic than that of dentrite is found in each coating, i.e. the decline of friction coefficient curve when the tip scratched across the eutectic. 3.3. Electrochemical measurements
Fig. 5. Vickers hardness along the distance from coating to substrate of coatings H1 and H2.
3.3.1. Polarization curves The polarization curves of substrate and coatings H1, H2 are plotted in Fig. 9. The results indicate that the corrosion resistance
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Fig. 7. Images of the nano-scratch grooves at substrate, coatings H1 and H2.
Fig. 9. Polarization curves of substrate, coatings H1 and H2.
Fig. 8. Variation of lateral force and friction coefficient with sliding distance.
of both coatings is much better than that of substrate. The most important reason is the different chemical compositions, i.e. more than 98% Fe in steel substrate makes it easily corroded. Instead, for the Hastelloy C22 coatings, the dense Cr2O3 layer is formed on the coating surface, preventing the coating from corrosion. Moreover, coating H2 displays better corrosion resistance with higher corrosion potential (Ecorr) of 0.286 V and lower corrosion current density (Icorr) of 2.45 107 A/cm2, than coating H1 with the values of 0.301 V and 3.40 107 A/cm2. Coating H2 has more compact microstructure and smaller grain size than coating H1 as stated in Fig. 1, thus there are more grain boundary and dislocations in coating H2, which provide more active nucleation sites for quick formation of continuous and protective passive film. The passive film plays a key role in preventing Ni ions or electrons from migrating towards the surface to enhance electrochemical reaction and aggravate corrosion. There was very little weight loss of coatings H1 and H2 with shining surface after immersing in 3.5 wt.% NaCl solution for several months. The fact indicates that the anticorrosion property of Hastelloy C22 coating prepared by laser cladding is as good as that of bulk one. The weight loss rate related to corrosion current density (Icorr) can be roughly predicted through the assumption of uniform corrosion and Faraday’s law [28]:
Q ¼ DW n N A =M
ð3Þ
DW ¼ W=At ¼ iM=nF
ð4Þ
Table 3 Corrosion results of the substrate, Hastelloy C22 coatings H1 and H2. Samples
H1 coating
H2 coating
Substrate
Average atomic mass (g/mol) Icorr (A/cm2) DW (mg/cm2 s) P (mm/y)
58.33 3.40 107 6.85 108 0.0025
60.43 2.45 10-7 5.11 108 0.0019
56 1.45 105 4.31 106 0.1792
where Q = total charges (Coulomb), DW = weight loss rate (g/cm2 s), M = atomic mass of the metal (g/mol), n = number of electron/atoms consumed (around 3 for both coatings, 2 for substrate), NA = Avogadro constant (6.022 1023), A = exposure area (10 10 mm2 in this paper), t = corrosion time (s), i = current density (A/cm2), F = Faraday’s constant (96,485 Coulomb/mol). Furthermore, through the relation below, the change of weight loss rate can also be converted to the average penetration rate:
P ¼ 3:65DW=q
ð5Þ
where P is the average penetration rate (mm/y), q is coating density (8.69 g/cm3 for both coatings). The final corrosion results are given in Table 3, which illustrates that coating H2 demonstrates higher corrosion resistance than coating H1 and substrate. 3.3.2. EIS measurements Electrochemical impedance spectroscopy (EIS) was measured to evaluate the variation of electrochemical reactions in the substrate and coatings. The Nyquist plots of the substrate during 7 days and
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Fig. 10. EIS of substrate, coatings H1 and H2.
coatings H1 and H2 during 16 days in 3.5 wt.% NaCl solution are shown in Fig. 10. The substrate exhibits one capacitive semicircle in the whole frequency range, and then the semicircle turns small and depresses with the increase of immersing time. This reveals that the local corrosion might have occurred on the substrate. Comparatively, the Nyquist plots of coatings show the continuously increasing a quarter capacitive circle with immersing time and then tend to stabilize, indicating the gradual formation of stable passive film. Besides, larger capacitive quarter-circle radius for coating H2 than for coating H1 is found, which reveals more compact and thicker passive film in coating H2. 4. Conclusions The Hastelloy C22 coatings H1 and H2 with few defects were produced by laser cladding with scanning speeds of 6 and 12 mm/s, respectively. Planar, cellular and dendritic solidifications were observed on the cross-sections of both coatings. Coating H2 shows more compact microstructure than that of coating H1 due to the higher laser scanning speed. Higher micro-hardness was found in coating H2 than H1 due to the existence of large quantity of eutectic containing Mo6Ni6C hard phase and the compact microstructure. For the same reason, coating H2 exhibits better tribological property. The corrosion resistance of coating H2, characterized with high corrosion potential of 0.286 V and low corrosion current density of 2.45 107 A/cm2, is better than H1 with the values of 0.301 V and 3.40 107 A/cm2 due to the quick formation of thick passive film. The results reveal that the microstructures, mechanical properties and corrosion resistance of Hastelloy C22 coatings are influenced by laser scanning speed. References [1] A. Zocco, A. Perrone, M.F. Vignolo, S. Duhalde, I. Avram, C. Morales, T. Pérez, Appl. Surf. Sci. 208–209 (2003) 669–675.
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