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Laser surface engineered TiC coating on 6061 Al alloy: microstructure and wear
Applied Surface Science 153 Ž2000. 65–78 www.elsevier.nlrlocaterapsusc
Laser surface engineered TiC coating on 6061 Al alloy: microstructure and wear Lalitha R. Katipelli, Arvind Agarwal, Narendra B. Dahotre
)
Department of Materials Science and Engineering, Center for Laser Applications, UniÕersity of Tennessee Space Institute, Tullahoma TN 37388, USA Received 23 July 1999; accepted 2 September 1999
Abstract Hard and refractory TiC has been deposited on 6061 Al alloy by Laser Surface Engineering ŽLSE.. A ‘‘composite’’ coating is obtained with TiC particles of various shapes and sizes embedded in Al alloy–Ti matrix. The coating is uniform, continuous and free of cracks. The various reactions occurring during laser processing were thermodynamically analyzed and related to the experimental observations. Microhardness measurements suggested high hardness values in the coating region and a strong bonding at the coatingrsubstrate interface. Dry sliding wear tests were performed to measure the wear resistance and the coefficient of friction of the coating. Wear resistance of the coated surface was found to be high when compared to the substrate side. The coefficient of friction was found to be 0.64. q 2000 Elsevier Science B.V. All rights reserved. Keywords: Laser surface engineering; TiC particles; Wear; Coating; Bonding; Interface
1. Introduction Aluminum alloys are used in many applications due to their excellent properties such as high strength-to-weight ratio, good ductility, lightweight, availability and low cost w1x. However, the surface properties of Al alloys the hardness and wear resistances in particular are insufficient to fulfil many industrial requirements. Surface coating with ceramics can be a promising approach to this problem w2x. Hard TiC ceramics are well known for combining a
) Corresponding author. Tel.: q1-931-393-7495; fax: q1-931454-2271. E-mail address: [email protected] ŽN.B. Dahotre.
number of special properties that have made them of particular interest for a wide variety of applications. They are used as a wear-resistant coating for cutting tools and inserts and as diffusion barriers in semiconductor technology w3x. TiC exhibits a very high melting point and thermal stability, high hardness and excellent wear resistance, low coefficient of friction, and high electrical and thermal conductivities. Because of its high melting point, TiC is a promising material to be used as first-wall material in fusion reactors w3,4x. There are several processes of depositing ceramics on Al alloys such as flame spraying, plasma spraying, screen printing, electroplated coating, PVD, CVD, etc. w5x. These methods are not widely used, as they do not offer metallurgical bonding to the base
0169-4332r00r$ - see front matter q 2000 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 9 - 4 3 3 2 Ž 9 9 . 0 0 3 6 8 - 2
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material. In such case, Laser Surface Engineering ŽLSE. offers to improve the surface properties of Al alloys while keeping the bulk properties more or less intact. By means of a high power laser, very high temperatures can be reached to melt both the metal substrate and, sometimes, the ceramic particles. This may promote chemical reaction and wetting between ceramic and metal, and, as a result, may lead to a strongly bonded ceramic–metal interface after resolidification w2x. Moreover, LSE involves high cooling rates Ž10 3 –10 8 Krs. which produce meta-stable phases, leading to the development of a wide variety of microstructures with novel properties that cannot be produced by any conventional processing technique w6x. Many researchers have studied the effect of ceramic coating on Al alloys w7–11x. However, studies on the influence of laser treatment of Al alloys for ceramic coating are very limited w2,12,13x. In the present study, hard TiC is deposited on Al 6061 alloy by LSE technique.
2. Experimental procedure 2.1. Materials
2.3. Characterization
Commercially available TiC powder Ž99.5% purity and average powder size - 15 mm. and 6061 Al alloy were used in the present study. The TiC powder was supplied by CERAC Milwaukee, WI. The chemical composition of 6061 Al alloy is given in Table 1. 2.2. Coating process Plates of 6061 Al alloy with dimensions 12 = 12 = 1 in3 were cleaned using sand blasting. The powder precursor made of 90 wt.% TiC q 10 wt.% Si suspended in a 10-wt.% hydroxyl methyl cellulose water-based organic binder Žproprietary formulation
Table 1 Chemical composition of 6061 Al alloy in wt.% Mg
Si
made from commercially available resins used in the paint industry. was spray deposited on the 6061 Al substrate. The addition of 10 wt.% Si to the TiC powder was intended to increase the wettability and fluidity of Al w6x. The average sprayed precursor thickness was 150 mm. Sprayed coupons were dried at 708C for 1 h prior to laser processing. A 2-kW Rofin Sinar continuous wave Nd:YAG laser equipped with a fiber optic beam delivery system was employed for laser treatment of the sprayed substrates. The optical fiber was 17 m long and 600 mm in diameter. The laser beam was focused at 0.5 mm above the surface of the substrate. The lenses within the output-coupling module of fiber optic were configured to provide a beam of 3.5 mm wide line in spatial distribution onto the sample surface. Such configuration provides a rapid processing speed and reduces the overlap between the laser passes. There was a 20% overlap between two consecutive laser tracks. The laser beam power and traverse speed were maintained constant at 1.8 kW and 120 cmrmin, respectively. The coated samples were additionally cooled by mounting them on a water-cooled copper plate during laser processing.
Cu
Cr
Mn
Fe
Al
0.8–1.2 0.4–0.8 0.05–0.4 0.04–0.35 - 0.15 - 0.7 Bal.
Phase identification was carried out on a Philips ˚ Norelco X-ray diffractometer with CuK a Ž1.54 A wavelength. radiation, operating at 40 kV and 20 mA. X-ray diffraction ŽXRD. analysis was performed on both the worn and unworn surfaces of the coating to determine the possible changes in phases as a result of wear tests. Microstructural characterization was performed on an ISI super III-A scanning electron microscope ŽSEM.. The samples for metallography were prepared by sectioning the coated plate perpendicular to the laser tracks. The sectioned faces were polished on a Buehler Isomet 2000 cloth to a diamond finish. Keller’s reagent was used as the etchant. Energy dispersive spectroscopy ŽEDS. and X-ray mapping equipped with SEM were utilized to determine the elemental distribution in a semiquantitative manner. Macroscopic observations of the cross-section of the coating were also conducted under BX60M Olympus microscope. This high-resolution optical microscope was equipped with differ-
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ent modes of light to observe the cross-section in desired light conditions. Microhardness tests were performed on a Buehler Isomet microhardness tester with a Knoop indentor under a normal load of 200 g applied for 15 s. Dry sliding wear tests were carried out on a Block-on-Disc tribometer to determine both the weight loss and coefficient of friction with respect to time. The speed of the rotating disc was maintained at 4.4 mrs and a normal load of 4 lb was applied to the sample while sliding. These tests were carried out on coated coupons of size 20 = 25 mm2 for a total 20-min duration. Weight loss measurements were made after successive 2 min. Coefficient of friction was also computed simultaneously by an interfaced computer that acquired data in the form of voltage and current of the motor as a function of time.
generation of stoichiometric andror non-stoichiometric derivations. Primarily, in the system under investigation, the following reactions are possible. The free energy of formation and their temperature ranges are obtained from various sources w14–18x, Ti q C
3.1. Thermodynamic predictions
™ TiC
™ SiC
Ž 2.
DG s y71.258 q 0.0078T kJrmole Ž0 - T - 1700 K., DG s y120.3 q 0.0368T kJrmole Ž1700 - T 3200 K. w16x; 4r3 Al q C
It is essential to study the thermodynamics of the system to understand the nature of chemical reactions taking place during the coating process. Therefore, in the present study, attempts are made firstly to thermodynamically predict the possible reactions and, secondly, to verify the existence of these reaction products using analytical techniques such as EDS and XRD. Such combined thermodynamic and analytical approach provides an insight into the mechanical, chemical and thermophysical properties of the coated sample. However, it is pointed out that often the thermodynamic conditions that prevail during laser processing deviate from equilibrium conditions, thereby making equilibrium phase diagram consideration less applicable and more difficult to predict the process products. Due to the fact that laser processing is a very complex non-equilibrium process and also because of the lack of data in the open literature on various reaction products during such non-equilibrium processes, a first approximation that generation of products based on equilibrium phase diagrams can be considered. These equilibrium products can be considered as sources for further
Ž 1.
DG s y182.9 q 0.01T kJrmole Ž298 - T - 115 K., DG s y186.4 q 0.013T kJrmole Ž1155 - T - 2000 K. w14x. From the Ti–C phase diagram, it has been corroborated that TiC x Ž x F 1. exists as a single homogeneous carbide phase with a wide range of stoichiometry w15x. The reaction is reversible and, hence, it is possible for Ti and C dissociated from TiC to combine with other reacting elements. Accordingly, the following reactions may be considered: Si q C
3. Results and discussion
67
™ 1r3 Al C 4
3
Ž 3.
DG s y71.315 q 0.013T kJrmole Ž0 - T - 900 K., DG s y91.055 q 0.033T kJrmole Ž900 - T 3200 K. w16x; Ti q 3Al
™ TiAl
3
Ž 4.
DG s y52.503 q 0.021T kJrmole Ž0 - T - 3200 K. w17,18x; Ti q 2Si
™ TiSi
2
Ž 5.
DG s y134.19 q 0.0067T kJrmole Ž298.15 - T 1700 K., DG s y1.5738 q 1.0063T kJrmole Ž1700 - T - 1813 K. w19x. With the obtained values of free energy of formation, a graph of DG vs. T is plotted in Fig. 2. From the figure, it is clearly observed that TiC is the most stable compound followed by TiSi 2 . This fact is also evident from the XRD peaks of TiC and TiSi 2 shown in Fig. 1. As mentioned earlier, titanium carbide exists as a single homogeneous phase, TiC x with a wide homogeneity range. The crystal nature of such phase is cubic of NaCl type and for TiC of stoichiometric
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Fig. 1. XRD spectrum of the TiC-coated 6061 Al sample.
composition, the lattice parameter varies between ˚ Similarly, silicon carbide exists 4.315 and 4.324 A. in the form of cubic b-SiC and hexagonal a-SiC. b-SiC with cubic crystal structure of ZnS type has a ˚ w20x. The values of lattice parameter of 4.358 A lattice parameters for both TiC and SiC being close, the X-ray reflection peaks of these phases overlap as observed in Fig. 1. Even though thermodynamically Al 4 C 3 is expected to exist up to 29278C ŽFig. 2. in vacuum, it is stable up to 12008C and, at 22008C, it sublimes without melting w21x. In addition, the solubility of carbon in aluminum is only 0.02–0.04 wt.% at 1300–15008C and, at 1000–11008C, it is practically zero w21x. Also, the possibility of formation of Al 4 C 3 through carbothermic reduction of SiC exists. Thus, Al 4 C 3 may form in the reaction w22x 4Al q 3SiC
™ Al C q 3Si 4
3
Ž 6.
However, the content of SiC and the duration of contact between SiC and the molten Al controls the formation of Al 4 C 3 w23x and, for complete stability
of Al 4 C 3 , Si levels above 8 wt.% are required. In the present experiments, though Si in the precursor was 10 wt.%, it was not sure how much was available as free Si in the melt during processing. In view of all these scenarios, it is less likely that Al 4 C 3 will form. This was further confirmed by XRD analysis ŽFig. 1.. In addition to reaction Ž4., TiAl 3 can form via the following reaction w7x: TiC Ž s . q 3Al Ž s .
™ TiAl Ž s. q C 3
Ž 7.
In the temperature range of 1150–1800 K, the Gibbs energy of formation Ž DG . for this reaction is positive, which indicates that TiAl 3 is unstable for an ideally dilute carbon concentration of 1 wt.%. However, TiAl 3 is stable in the presence of TiC when DG becomes negative at concentration of dissolved C lower than the 10y5 wt.% w17x. Therefore, TiAl 3 formation can only be expected during processing if Ti concentration is very high. Such extremely low concentration of C may be possible during the present process due to formation of SiC via reaction Ž2..
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Fig. 2. Gibbs free energy of formation of different compounds that may possibly form during laser processing.
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the metals. The better wetting of refractory carbides by transition elements is attributed to the possible capture of non-localized valence electrons of molten metal atoms by the disturbed configuration of carbon atoms in the carbides w25x. Table 2 lists the wetting angles of certain refractory carbides by molten metals. The wetting between TiC and Al seems to have improved due to the addition of Si in the precursor and Mg present in 6061 Al substrate ŽTable 1., which is evident from the strong and adherent coating of the sample. Table 2 shows the low contact angle values suggesting high wetting of TiC by molten Si and Mg. It has been earlier observed that the addition of Mg also leads to the formation of spinel at the metalrceramic interface, which enhances the wetting w26x. In the present study, formation of spinel formation was not observed by XRD. However, a detailed investigation by transmission electron microscope may reveal localized spinel formation that promotes wetting. A thick laser melt zone ŽLMZ. comprising of the fine dendrites of Al alloy is formed in the substrate underneath the coating. In LSE, the rapid solidification process can result in a considerable supercooling resulting in the refinement of microstructure from coarse to fine dendritic structure in the melt zone w6x.
3.2. Microstructural characterization Fig. 3 shows an overview of the cross-section of the TiC-coated 6061 Al sample. The coating is dense and adherent to the substrate. The thickness of the coating is uniform and was found to be approximately 125 mm. The reduction of about 20% from 150 mm thickness of the precursor deposit to the final coating thickness of 125 mm could be attributed primarily to evaporation of the entire binder material along with possible evaporation of a small amount of the precursor powder mixture material. The evaporation of these materials is also expected to reduce the porosity between the powder particles, thereby providing a dense and less thick coating. The coating is composite in nature with TiC particles embedded in the Al alloy matrix. It has been experimentally proven that refractory carbides are practically not wetted by Group III elements and are, in general, wetted by transition elements w24x. Such difference in the wetting nature is explained by the electronic structure of
Fig. 3. SEM micrograph of the overview of the cross-section of the TiC-coated 6061 Al sample.
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Table 2 Wetting of some refractory carbides by molten metals Carbide
Wetting melt
Temperature Ž8C.
Atmosphere
Contact angle in degrees
Reference
TiC
Fe Co Ni Si Mg Al Fe Co Ni Si Mg Al
1550 1500 1450 1500 1300 700 1500 1500 1500 1500 Not available 1000
Hydrogen Hydrogen Hydrogen Vacuum Vacuum Vacuum Vacuum Hydrogen Vacuum Hydrogen Not available Vacuum
39 36 17 32 50 118 0 0 0 0 Not available 135
w24x w24x w24x w14x w14x w14x w24x w24x w24x w14x Not available w14x
WC
The LMZ has an approximate depth of 300 mm. Fig. 4 shows the high magnification micrograph of the interface of LMZ with the 6061 Al substrate. The interfacial region shows the growth of dendrites from the LMZ towards the coating. Fig. 5Ža. shows the SEM micrograph of the cross-section of the coating. TiC particles of various size and shape within the matrix of 6061 Al are observed. The volume fraction of TiC particles in the coating is found to be approximately 65%. The elemental X-ray maps of the coating corresponding to Ti, Si and Al distribution are shown in Fig. 5Žb.,
Fig. 4. SEM micrograph of LMZ-substrate interface.
Žc., and Žd., respectively. Ti-rich zone shows the presence of TiC particles whereas the coating matrix is distinguished by the Al rich zone. Si is distributed uniformly over the entire surface. Table 3 provides the results of the EDS quantitative analysis conducted at various locations in the coating. These results suggest that the matrix within the coating contains Al Ž53.24 at.%. and Ti Ž35.86 at.%. along with some Si Ž10.88 at.%.. Such high content of Ti in the matrix could possibly be attributed to fragmentation andror dissolution of TiC particles. Fragmentation of such TiC particles is observed in both Figs. 3 and 5Ža.. Partial dissolution of TiC and, thereby, availability of free Ti in the matrix is desirable as Ti, being a highly reactive element, tends to modify the surface properties of the carbide particles for enhanced wettability with the substrate molten material. Similar is the case with the interface between TiC particles and the coating matrix, which contains Ti Ž35.19 at.%. along with Al Ž59.053 at.%. and a trace amount of Si Ž5.763 at.%.. Such nature of the reaction product is expected to provide chemical bonding in addition to mechanical bonding between the TiC particles and the matrix. Fig. 6Ža. shows the SEM micrograph of the coating-substrate interface. The corresponding X-ray elemental distribution of Al, Ti and Si is shown in Fig. 6Žb., Žc. and Žd., respectively. Due to a large difference in the coefficient of the thermal expansions of metals and ceramics, high residual stresses may occur at the interface during the rapid solidification after laser processing. These residual stresses can result in the delaminationrbuckling of the coating
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Fig. 5. Ža. High magnification SEM micrograph of the cross-section of the coating and corresponding X-ray elemental map of Žb. Al, Žc. Si, and Žd. Ti.
from the substrate w2x. However, in the present case, no such delamination was observed. This indicates that the coating is sound and adherent to the substrate suggesting a metallurgical bond. A sharp interface in all these figures suggests no transferrdiffusion of Ti into the region of the substrate material surrounding the interface. On the contrary, Al has flown into the matrix of the coating ŽFig. 6b.. Table 3 EDS analysis in the coating Location
Matrix within the coating TiC particles Interface of TiC particles and coating matrix
At.% Ti
Al
Si
35.86 96.589 35.19
53.24 0.484 59.053
10.88 2.927 5.763
Also, interfacial adhesion strength between ceramic particle and metallic melt is often represented by the work of adhesion, Wa , which is defined as work per unit area of interface, necessary to separate reversibly a solid–liquid interface to create a solid– vapor interface w27x. Work of adhesion, Wa , for several combination of w28x metalrceramic system is tabulated in Table 4. The work of adhesion, Wa , reflects directly the importance of energetical interactions between the solid and liquid phases. A higher value Wa suggests stronger interactions w27x. Wa for TiCrSi and TiCrMg is significantly higher than TiCrAl ŽTable 4.. Thus, in the present study, the presence of Si and Mg in the precursor and 6061 Al is expected to increase the wettability and, hence, interfacial strength between TiC and Al melt, thereby resulting in an adherent coating.
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Fig. 6. Ža. SEM micrograph of the coating–substrate interface and corresponding X-ray elemental map of Žb. Al, Žc. Si and Žd. Ti.
3.3. Mechanical characterization Mechanical features of the TiC coated 6061 Al alloy were studied using a microhardness tester. A
SEM micrograph of the cross-section of the coating–substrate interface with microhardness indentations is shown in Fig. 7Ža.. X-ray elemental distribution of Al, Si, and Ti corresponding to
Table 4 Work of adhesion for metalrceramic systems Ceramic
Metallic melt
Surface energy, g 1y Õ of liquid–vapor interface ŽmJrm2 .
Contact angle in degrees
Work of adhesion, Wa ŽmJrm2 .
TiC
Al Si Mg Fe Al Si Mg Fe
914 860 583 1900 915 860 Not available 1785
118 32 50 39 160 34 Not available 49
485 1720 960 3578 55 1575 Not available 2955
TiB 2
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Fig. 7. Ža. SEM micrograph of the coating, coating–substrate interface, and substrate regions with indentations and corresponding X-ray elemental map of Žb. Al, Žc. Si, and Žd. Ti.
Fig. 7Ža. are represented in Fig. 7Žb., Žc. and Žd., respectively. The relative difference in size of microhardness indentations explains the variation in hardness. Considerably high microhardness values were observed in the coating. The two indentations taken in the coating showed different values of microhardness ŽKnoop., with the top part of the coating Ž572 " 60. relatively harder than the bottom part Ž454 " 30.. This is due to the presence of SiC, which apparently segregated to the top of the coating during laser processing. SiC, being less denser Ž3.1 to 3.2 grcm3 . in comparison to TiC Ž4.92 grcm3 ., floats to the top of molten coating during laser processing and remains at the top of the coating after resolidification. This fact is also evident from the corresponding X-ray elemental map of Si shown in Fig. 7Žc.. The hardness of the LMZ is 72 " 5 and that of unreacted substrate is 84 " 10. The relatively
low hardness value of LMZ is attributed to the dissolutionrredistribution of precipitates formed during the precipitation hardening of 6061 Al alloy. In the present study, attempts were made to semiquantitatively determine the interfacial strength. In the process, microhardness indentations were made
Table 5 SEM analysis of the indentations at the interface Applied load Žg.
Length of the longer diagonal Žmm.
SEM predictions
50 100 200 300 500 1000
32 61 72 107 116 150
No interfacial cracks No interfacial cracks No interfacial cracks No interfacial cracks No interfacial cracks No interfacial cracks
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along the interface at various loads Ž50, 100, 200, 300, 500, and 1000 g. to check for initiation of any cracks or delaminations at the site of indentations. The loads at which the microhardness tests are conducted and the corresponding SEM observations are illustrated in Table 5. Fig. 8 shows the corresponding
SEM micrographs of the indentations taken at all the test loads within the limitations of the indentor being used in the present study. SEM observations showed that there were no visible cracks or delaminations at these indentations. However, the possibility of fine stationary cracks within the coating is not ruled out
Fig. 8. SEM micrographs of the indentations taken at the coating–substrate interface at various loads.
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since this can only be observed at very high magnification. The dynamic behavior of these stationary cracks is being further investigated using a four-point bending test. Such tests will characterize the strength of ceramic–metal interface in a quantitative manner. 3.4. Tribological characterization 3.4.1. Wear tests The tribological performance of a ceramic coating depends on the characteristics such as hardness, thickness, internal stress level, and load-bearing capacity. When a coating cools from the high temperature of the melt zone in the deposition process, the internal compressive stress developed within the coating reduces the extent of surface fracture and, thus, increases the sliding wear resistance of the coated surface w29x. In addition, wear behavior of any material system is expected to be influenced by environment, such as dry Žair or inert atmosphere. and wet Žlubricated.. The thermo–physical Žforce and temperature. conditions at the point of contact are extreme during wear testing. These extreme conditions are favorable to cause degradationrdissociation Žphysical and chemical. of the material and, hence, occurrence of localized reactionŽs. between the material and surrounding environment, thereby providing different wear responses. Wear results in the loss of material, elevation of surface temperature, surface activation, and the removal of protective film, and, thus, may accelerate oxidation or corrosion w30x. Hence, it is important to
Fig. 9. Weight loss vs. time for the wear tests.
75
study the wear behavior of the coated samples to determine the nature of the coating. Fig. 9 shows a plot of weight loss vs. time for the wear tests conducted on TiC coated and uncoated Žsubstrate. surfaces of the laser-treated samples. It is clearly observed that the weight loss for the coated side was lower than the substrate side. The wear rate is also significantly improved. The total duration of the wear test was 20 min. However, while conducting the tests on the uncoated substrate side, the substantial weight losses were calculated for a time duration of only 30 s due to high seizure of the substrate material during sliding. Comparison of XRD analyses of the worn and unworn-coated surfaces as illustrated in Fig. 1 indicated that there are no new phases formed during the wear process. This suggested that the coating was chemically and physically stable under the dry sliding conditions used in the present study. Surface roughness profiles shown in Fig. 10Ža., Žb. and Žc. indicated that the average surface roughness, R a , of the test sample increased from 0.43 to 13.53 mm after laser processing. The average surface roughness of the worn surface is 6.51 mm. Topographic features of the worn surface were analyzed using SEM. Fig. 11Ža. shows a low magnification micrograph of the TiC coated 6061 Al sample. X-ray elemental distributions of Al, Ti, and Fe corresponding to Fig. 11Ža. are shown in Fig. 11Žb., Žc. and Žd., respectively. The worn surface of the coating is identified by continuous wear scars parallel to the laser tracks in some isolated regions of contact between the sample and wear-disc. During sliding, there is a material-to-material contact, followed by welding or fusing of the contacting asperities. When sliding continues, the asperities of the weaker material may shear off and transfer to the opposite surface w31x. In the present work, the hardened steel sliding disc Ž62 Rc., whose main constituent is Fe, shears off the Al present as the matrix of the coating, and Al is collected as the loose debris at the end of the scars. On the contrary, TiC present in the coating, being harder than Fe, shears off the material from the sliding disc, and, as a result, Fe is deposited in those scars. Such composite nature of wear and segregation of Al and Fe in particular regions of the worn surface are clearly seen from Fig. 11Ža., Žb. and Žd., respectively.
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Fig. 10. Optical profilometer figures showing the surface roughness of Ža. 6061 Al substrate, Žb. coated side and Žc. worn coated surface.
3.4.2. Coefficient of friction measurement The measurement of the coefficient of friction provides direct information about the work done to deform the surface of the material. However, it is not necessarily a direct indication of material loss or
separation of loose debris from the surface w15x. In the present study, the tribometer interfaced with a computer recorded the wear test parameters such as Voltage Ž V . and Current Ž I . as a function of time. Based on the principle of energy conservation, the
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Fig. 11. Ža. SEM micrograph of the worn surface and corresponding X-ray elemental map of Žb. Al, Žc. Ti and Žd. Fe.
frictional force is equal to the electrical work done by the motor given by
The frictional theory states that Wf s m NÕ
Wf s Voltage Ž DV . = Current Ž D I .
where N is the normal load, Õ is the linear speed of the sliding disc and m is the coefficient of friction. By equating the above equations, the coefficient of friction can be computed. The computed coefficient of friction for TiC coating has been plotted for the entire test time of 20 min as shown in Fig. 12. The best-fit line shows the coefficient of friction of the TiC-coated surface to be approximately 0.64.
Ž 8.
Ž 9.
4. Conclusions
Fig. 12. Coefficient of friction vs. time for the wear tests.
1. TiC has been deposited on 6061 Al alloy using the LSE technique. The coating is uniform, adherent, and free of cracks and porosities.
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2. The coating is composite in nature with the microstructure consisting primarily of hard TiC particles reinforced in a matrix of Al–Ti mixture. 3. The average surface roughness of the coating was found to be 13.5 mm. 4. Considerably high hardness values are obtained in the coating. 5. The coating is wear resistant and the wear resistance is highly influenced by the composite nature of the coating. 6. The coefficient of friction was computed to be approximately 0.64.
Acknowledgements Authors acknowledge the partial support from the Surface Technology Division of ALCOA.
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w10x H.W. Bergmann, Th. Endres, B. Juckenath, LIA 69 Ž1989. 239, ICALEO. w11x D. Pantelis, E. Giannetaki, Y. Chryssoulakis, P. Ponthiaux, Plat. Surf. Eng. Ž1994. 52. w12x X.B. Zhou, J.Th.M. De Hosson, Acta Metall. Mater. 42 Ž4. Ž1994. 1155. w13x H.J. Hegge, J. Boetje, J.Th.M. De Hosson, J. Mater. Sci. 25 Ž1990. 2335. w14x G.M. Samsonov, I.M. Vinitskii, in: Handbook of Refractory Compounds, IFI Plenum, New York, 1980, p. 217. w15x A. Agarwal, N.B. Dahotre, J. Mater. Eng. Perform. 8 Ž4. Ž1999. 479. w16x JANAF Thermodynamic Tables, 3rd edn. Part-I, Al–Co. w17x R.A. Rapp, X. Zhebg, Metall. Trans. A 22A Ž1991. 3071. w18x J.L. Murray, Metall. Trans. A 19A Ž1988. 243. w19x T.Ya. Kosolapova, in: Handbook of High Temperature Compounds: Properties, Production and Applications, Hemisphere Publishing, New York, 1990, p. 252. w20x T.Ya. Kosolapova, in: Carbides: Properties, Production and Applications, Plenum, New York, 1971, p. 16. w21x T.Ya. Kosolapova, in: Carbides: Properties, Production and Applications, Plenum, New York, 1971, p. 88. w22x N.B. Dahotre, M.H. McCay, T.D. McCay, S. Gopinathan, L.F. Allard, , J. Mater. Res. 6 Ž3. Ž1991. 514. w23x N.B. Dahotre, T.D. McCay, M.H. McCay, J. Appl. Phys. 65 Ž12. Ž1989. 5072. w24x G.S. Upadhyaya, in: Sintered Metal–Ceramics Composites, Elsevier, Amsterdam, 1984, p. 41. w25x G.S. Upadhyaya, in: Synthetic Materials for Electronics, Elsevier, Amsterdam, 1981, p. 133. w26x X.B. Zhou, P.M. Bronsvled, J.Th.M. De Hosson, Lasers Eng. 1 Ž1991. 145. w27x J.G. Li, Ceram. Int. 20 Ž1994. 391. w28x T.Ya. Kosolapova, in: Handbook of High Temperature Compounds: Properties, Production and Applications, Hemisphere Publishing, New York, 1990, p. 685. w29x S. Bahadur, C.-N. Yang, Wear 196 Ž1996. 156. w30x C.K. Fang, C.C. Huang, T.H. Chuang, Metall. Trans. A 30A Ž1999. 643, ŽMarch.. w31x J. Kelly, K. Nagarathnam, J. Majumder, J. Laser Appl. 10 Ž2. Ž1998. 45, ŽApril..
Laser surface engineered TiC coating on 6061 Al alloy: microstructure and wear Lalitha R. Katipelli, Arvind Agarwal, Narendra B. Dahotre
)
Department of Materials Science and Engineering, Center for Laser Applications, UniÕersity of Tennessee Space Institute, Tullahoma TN 37388, USA Received 23 July 1999; accepted 2 September 1999
Abstract Hard and refractory TiC has been deposited on 6061 Al alloy by Laser Surface Engineering ŽLSE.. A ‘‘composite’’ coating is obtained with TiC particles of various shapes and sizes embedded in Al alloy–Ti matrix. The coating is uniform, continuous and free of cracks. The various reactions occurring during laser processing were thermodynamically analyzed and related to the experimental observations. Microhardness measurements suggested high hardness values in the coating region and a strong bonding at the coatingrsubstrate interface. Dry sliding wear tests were performed to measure the wear resistance and the coefficient of friction of the coating. Wear resistance of the coated surface was found to be high when compared to the substrate side. The coefficient of friction was found to be 0.64. q 2000 Elsevier Science B.V. All rights reserved. Keywords: Laser surface engineering; TiC particles; Wear; Coating; Bonding; Interface
1. Introduction Aluminum alloys are used in many applications due to their excellent properties such as high strength-to-weight ratio, good ductility, lightweight, availability and low cost w1x. However, the surface properties of Al alloys the hardness and wear resistances in particular are insufficient to fulfil many industrial requirements. Surface coating with ceramics can be a promising approach to this problem w2x. Hard TiC ceramics are well known for combining a
) Corresponding author. Tel.: q1-931-393-7495; fax: q1-931454-2271. E-mail address: [email protected] ŽN.B. Dahotre.
number of special properties that have made them of particular interest for a wide variety of applications. They are used as a wear-resistant coating for cutting tools and inserts and as diffusion barriers in semiconductor technology w3x. TiC exhibits a very high melting point and thermal stability, high hardness and excellent wear resistance, low coefficient of friction, and high electrical and thermal conductivities. Because of its high melting point, TiC is a promising material to be used as first-wall material in fusion reactors w3,4x. There are several processes of depositing ceramics on Al alloys such as flame spraying, plasma spraying, screen printing, electroplated coating, PVD, CVD, etc. w5x. These methods are not widely used, as they do not offer metallurgical bonding to the base
0169-4332r00r$ - see front matter q 2000 Elsevier Science B.V. All rights reserved. PII: S 0 1 6 9 - 4 3 3 2 Ž 9 9 . 0 0 3 6 8 - 2
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material. In such case, Laser Surface Engineering ŽLSE. offers to improve the surface properties of Al alloys while keeping the bulk properties more or less intact. By means of a high power laser, very high temperatures can be reached to melt both the metal substrate and, sometimes, the ceramic particles. This may promote chemical reaction and wetting between ceramic and metal, and, as a result, may lead to a strongly bonded ceramic–metal interface after resolidification w2x. Moreover, LSE involves high cooling rates Ž10 3 –10 8 Krs. which produce meta-stable phases, leading to the development of a wide variety of microstructures with novel properties that cannot be produced by any conventional processing technique w6x. Many researchers have studied the effect of ceramic coating on Al alloys w7–11x. However, studies on the influence of laser treatment of Al alloys for ceramic coating are very limited w2,12,13x. In the present study, hard TiC is deposited on Al 6061 alloy by LSE technique.
2. Experimental procedure 2.1. Materials
2.3. Characterization
Commercially available TiC powder Ž99.5% purity and average powder size - 15 mm. and 6061 Al alloy were used in the present study. The TiC powder was supplied by CERAC Milwaukee, WI. The chemical composition of 6061 Al alloy is given in Table 1. 2.2. Coating process Plates of 6061 Al alloy with dimensions 12 = 12 = 1 in3 were cleaned using sand blasting. The powder precursor made of 90 wt.% TiC q 10 wt.% Si suspended in a 10-wt.% hydroxyl methyl cellulose water-based organic binder Žproprietary formulation
Table 1 Chemical composition of 6061 Al alloy in wt.% Mg
Si
made from commercially available resins used in the paint industry. was spray deposited on the 6061 Al substrate. The addition of 10 wt.% Si to the TiC powder was intended to increase the wettability and fluidity of Al w6x. The average sprayed precursor thickness was 150 mm. Sprayed coupons were dried at 708C for 1 h prior to laser processing. A 2-kW Rofin Sinar continuous wave Nd:YAG laser equipped with a fiber optic beam delivery system was employed for laser treatment of the sprayed substrates. The optical fiber was 17 m long and 600 mm in diameter. The laser beam was focused at 0.5 mm above the surface of the substrate. The lenses within the output-coupling module of fiber optic were configured to provide a beam of 3.5 mm wide line in spatial distribution onto the sample surface. Such configuration provides a rapid processing speed and reduces the overlap between the laser passes. There was a 20% overlap between two consecutive laser tracks. The laser beam power and traverse speed were maintained constant at 1.8 kW and 120 cmrmin, respectively. The coated samples were additionally cooled by mounting them on a water-cooled copper plate during laser processing.
Cu
Cr
Mn
Fe
Al
0.8–1.2 0.4–0.8 0.05–0.4 0.04–0.35 - 0.15 - 0.7 Bal.
Phase identification was carried out on a Philips ˚ Norelco X-ray diffractometer with CuK a Ž1.54 A wavelength. radiation, operating at 40 kV and 20 mA. X-ray diffraction ŽXRD. analysis was performed on both the worn and unworn surfaces of the coating to determine the possible changes in phases as a result of wear tests. Microstructural characterization was performed on an ISI super III-A scanning electron microscope ŽSEM.. The samples for metallography were prepared by sectioning the coated plate perpendicular to the laser tracks. The sectioned faces were polished on a Buehler Isomet 2000 cloth to a diamond finish. Keller’s reagent was used as the etchant. Energy dispersive spectroscopy ŽEDS. and X-ray mapping equipped with SEM were utilized to determine the elemental distribution in a semiquantitative manner. Macroscopic observations of the cross-section of the coating were also conducted under BX60M Olympus microscope. This high-resolution optical microscope was equipped with differ-
L.R. Katipelli et al.r Applied Surface Science 153 (2000) 65–78
ent modes of light to observe the cross-section in desired light conditions. Microhardness tests were performed on a Buehler Isomet microhardness tester with a Knoop indentor under a normal load of 200 g applied for 15 s. Dry sliding wear tests were carried out on a Block-on-Disc tribometer to determine both the weight loss and coefficient of friction with respect to time. The speed of the rotating disc was maintained at 4.4 mrs and a normal load of 4 lb was applied to the sample while sliding. These tests were carried out on coated coupons of size 20 = 25 mm2 for a total 20-min duration. Weight loss measurements were made after successive 2 min. Coefficient of friction was also computed simultaneously by an interfaced computer that acquired data in the form of voltage and current of the motor as a function of time.
generation of stoichiometric andror non-stoichiometric derivations. Primarily, in the system under investigation, the following reactions are possible. The free energy of formation and their temperature ranges are obtained from various sources w14–18x, Ti q C
3.1. Thermodynamic predictions
™ TiC
™ SiC
Ž 2.
DG s y71.258 q 0.0078T kJrmole Ž0 - T - 1700 K., DG s y120.3 q 0.0368T kJrmole Ž1700 - T 3200 K. w16x; 4r3 Al q C
It is essential to study the thermodynamics of the system to understand the nature of chemical reactions taking place during the coating process. Therefore, in the present study, attempts are made firstly to thermodynamically predict the possible reactions and, secondly, to verify the existence of these reaction products using analytical techniques such as EDS and XRD. Such combined thermodynamic and analytical approach provides an insight into the mechanical, chemical and thermophysical properties of the coated sample. However, it is pointed out that often the thermodynamic conditions that prevail during laser processing deviate from equilibrium conditions, thereby making equilibrium phase diagram consideration less applicable and more difficult to predict the process products. Due to the fact that laser processing is a very complex non-equilibrium process and also because of the lack of data in the open literature on various reaction products during such non-equilibrium processes, a first approximation that generation of products based on equilibrium phase diagrams can be considered. These equilibrium products can be considered as sources for further
Ž 1.
DG s y182.9 q 0.01T kJrmole Ž298 - T - 115 K., DG s y186.4 q 0.013T kJrmole Ž1155 - T - 2000 K. w14x. From the Ti–C phase diagram, it has been corroborated that TiC x Ž x F 1. exists as a single homogeneous carbide phase with a wide range of stoichiometry w15x. The reaction is reversible and, hence, it is possible for Ti and C dissociated from TiC to combine with other reacting elements. Accordingly, the following reactions may be considered: Si q C
3. Results and discussion
67
™ 1r3 Al C 4
3
Ž 3.
DG s y71.315 q 0.013T kJrmole Ž0 - T - 900 K., DG s y91.055 q 0.033T kJrmole Ž900 - T 3200 K. w16x; Ti q 3Al
™ TiAl
3
Ž 4.
DG s y52.503 q 0.021T kJrmole Ž0 - T - 3200 K. w17,18x; Ti q 2Si
™ TiSi
2
Ž 5.
DG s y134.19 q 0.0067T kJrmole Ž298.15 - T 1700 K., DG s y1.5738 q 1.0063T kJrmole Ž1700 - T - 1813 K. w19x. With the obtained values of free energy of formation, a graph of DG vs. T is plotted in Fig. 2. From the figure, it is clearly observed that TiC is the most stable compound followed by TiSi 2 . This fact is also evident from the XRD peaks of TiC and TiSi 2 shown in Fig. 1. As mentioned earlier, titanium carbide exists as a single homogeneous phase, TiC x with a wide homogeneity range. The crystal nature of such phase is cubic of NaCl type and for TiC of stoichiometric
L.R. Katipelli et al.r Applied Surface Science 153 (2000) 65–78
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Fig. 1. XRD spectrum of the TiC-coated 6061 Al sample.
composition, the lattice parameter varies between ˚ Similarly, silicon carbide exists 4.315 and 4.324 A. in the form of cubic b-SiC and hexagonal a-SiC. b-SiC with cubic crystal structure of ZnS type has a ˚ w20x. The values of lattice parameter of 4.358 A lattice parameters for both TiC and SiC being close, the X-ray reflection peaks of these phases overlap as observed in Fig. 1. Even though thermodynamically Al 4 C 3 is expected to exist up to 29278C ŽFig. 2. in vacuum, it is stable up to 12008C and, at 22008C, it sublimes without melting w21x. In addition, the solubility of carbon in aluminum is only 0.02–0.04 wt.% at 1300–15008C and, at 1000–11008C, it is practically zero w21x. Also, the possibility of formation of Al 4 C 3 through carbothermic reduction of SiC exists. Thus, Al 4 C 3 may form in the reaction w22x 4Al q 3SiC
™ Al C q 3Si 4
3
Ž 6.
However, the content of SiC and the duration of contact between SiC and the molten Al controls the formation of Al 4 C 3 w23x and, for complete stability
of Al 4 C 3 , Si levels above 8 wt.% are required. In the present experiments, though Si in the precursor was 10 wt.%, it was not sure how much was available as free Si in the melt during processing. In view of all these scenarios, it is less likely that Al 4 C 3 will form. This was further confirmed by XRD analysis ŽFig. 1.. In addition to reaction Ž4., TiAl 3 can form via the following reaction w7x: TiC Ž s . q 3Al Ž s .
™ TiAl Ž s. q C 3
Ž 7.
In the temperature range of 1150–1800 K, the Gibbs energy of formation Ž DG . for this reaction is positive, which indicates that TiAl 3 is unstable for an ideally dilute carbon concentration of 1 wt.%. However, TiAl 3 is stable in the presence of TiC when DG becomes negative at concentration of dissolved C lower than the 10y5 wt.% w17x. Therefore, TiAl 3 formation can only be expected during processing if Ti concentration is very high. Such extremely low concentration of C may be possible during the present process due to formation of SiC via reaction Ž2..
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Fig. 2. Gibbs free energy of formation of different compounds that may possibly form during laser processing.
69
the metals. The better wetting of refractory carbides by transition elements is attributed to the possible capture of non-localized valence electrons of molten metal atoms by the disturbed configuration of carbon atoms in the carbides w25x. Table 2 lists the wetting angles of certain refractory carbides by molten metals. The wetting between TiC and Al seems to have improved due to the addition of Si in the precursor and Mg present in 6061 Al substrate ŽTable 1., which is evident from the strong and adherent coating of the sample. Table 2 shows the low contact angle values suggesting high wetting of TiC by molten Si and Mg. It has been earlier observed that the addition of Mg also leads to the formation of spinel at the metalrceramic interface, which enhances the wetting w26x. In the present study, formation of spinel formation was not observed by XRD. However, a detailed investigation by transmission electron microscope may reveal localized spinel formation that promotes wetting. A thick laser melt zone ŽLMZ. comprising of the fine dendrites of Al alloy is formed in the substrate underneath the coating. In LSE, the rapid solidification process can result in a considerable supercooling resulting in the refinement of microstructure from coarse to fine dendritic structure in the melt zone w6x.
3.2. Microstructural characterization Fig. 3 shows an overview of the cross-section of the TiC-coated 6061 Al sample. The coating is dense and adherent to the substrate. The thickness of the coating is uniform and was found to be approximately 125 mm. The reduction of about 20% from 150 mm thickness of the precursor deposit to the final coating thickness of 125 mm could be attributed primarily to evaporation of the entire binder material along with possible evaporation of a small amount of the precursor powder mixture material. The evaporation of these materials is also expected to reduce the porosity between the powder particles, thereby providing a dense and less thick coating. The coating is composite in nature with TiC particles embedded in the Al alloy matrix. It has been experimentally proven that refractory carbides are practically not wetted by Group III elements and are, in general, wetted by transition elements w24x. Such difference in the wetting nature is explained by the electronic structure of
Fig. 3. SEM micrograph of the overview of the cross-section of the TiC-coated 6061 Al sample.
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Table 2 Wetting of some refractory carbides by molten metals Carbide
Wetting melt
Temperature Ž8C.
Atmosphere
Contact angle in degrees
Reference
TiC
Fe Co Ni Si Mg Al Fe Co Ni Si Mg Al
1550 1500 1450 1500 1300 700 1500 1500 1500 1500 Not available 1000
Hydrogen Hydrogen Hydrogen Vacuum Vacuum Vacuum Vacuum Hydrogen Vacuum Hydrogen Not available Vacuum
39 36 17 32 50 118 0 0 0 0 Not available 135
w24x w24x w24x w14x w14x w14x w24x w24x w24x w14x Not available w14x
WC
The LMZ has an approximate depth of 300 mm. Fig. 4 shows the high magnification micrograph of the interface of LMZ with the 6061 Al substrate. The interfacial region shows the growth of dendrites from the LMZ towards the coating. Fig. 5Ža. shows the SEM micrograph of the cross-section of the coating. TiC particles of various size and shape within the matrix of 6061 Al are observed. The volume fraction of TiC particles in the coating is found to be approximately 65%. The elemental X-ray maps of the coating corresponding to Ti, Si and Al distribution are shown in Fig. 5Žb.,
Fig. 4. SEM micrograph of LMZ-substrate interface.
Žc., and Žd., respectively. Ti-rich zone shows the presence of TiC particles whereas the coating matrix is distinguished by the Al rich zone. Si is distributed uniformly over the entire surface. Table 3 provides the results of the EDS quantitative analysis conducted at various locations in the coating. These results suggest that the matrix within the coating contains Al Ž53.24 at.%. and Ti Ž35.86 at.%. along with some Si Ž10.88 at.%.. Such high content of Ti in the matrix could possibly be attributed to fragmentation andror dissolution of TiC particles. Fragmentation of such TiC particles is observed in both Figs. 3 and 5Ža.. Partial dissolution of TiC and, thereby, availability of free Ti in the matrix is desirable as Ti, being a highly reactive element, tends to modify the surface properties of the carbide particles for enhanced wettability with the substrate molten material. Similar is the case with the interface between TiC particles and the coating matrix, which contains Ti Ž35.19 at.%. along with Al Ž59.053 at.%. and a trace amount of Si Ž5.763 at.%.. Such nature of the reaction product is expected to provide chemical bonding in addition to mechanical bonding between the TiC particles and the matrix. Fig. 6Ža. shows the SEM micrograph of the coating-substrate interface. The corresponding X-ray elemental distribution of Al, Ti and Si is shown in Fig. 6Žb., Žc. and Žd., respectively. Due to a large difference in the coefficient of the thermal expansions of metals and ceramics, high residual stresses may occur at the interface during the rapid solidification after laser processing. These residual stresses can result in the delaminationrbuckling of the coating
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Fig. 5. Ža. High magnification SEM micrograph of the cross-section of the coating and corresponding X-ray elemental map of Žb. Al, Žc. Si, and Žd. Ti.
from the substrate w2x. However, in the present case, no such delamination was observed. This indicates that the coating is sound and adherent to the substrate suggesting a metallurgical bond. A sharp interface in all these figures suggests no transferrdiffusion of Ti into the region of the substrate material surrounding the interface. On the contrary, Al has flown into the matrix of the coating ŽFig. 6b.. Table 3 EDS analysis in the coating Location
Matrix within the coating TiC particles Interface of TiC particles and coating matrix
At.% Ti
Al
Si
35.86 96.589 35.19
53.24 0.484 59.053
10.88 2.927 5.763
Also, interfacial adhesion strength between ceramic particle and metallic melt is often represented by the work of adhesion, Wa , which is defined as work per unit area of interface, necessary to separate reversibly a solid–liquid interface to create a solid– vapor interface w27x. Work of adhesion, Wa , for several combination of w28x metalrceramic system is tabulated in Table 4. The work of adhesion, Wa , reflects directly the importance of energetical interactions between the solid and liquid phases. A higher value Wa suggests stronger interactions w27x. Wa for TiCrSi and TiCrMg is significantly higher than TiCrAl ŽTable 4.. Thus, in the present study, the presence of Si and Mg in the precursor and 6061 Al is expected to increase the wettability and, hence, interfacial strength between TiC and Al melt, thereby resulting in an adherent coating.
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Fig. 6. Ža. SEM micrograph of the coating–substrate interface and corresponding X-ray elemental map of Žb. Al, Žc. Si and Žd. Ti.
3.3. Mechanical characterization Mechanical features of the TiC coated 6061 Al alloy were studied using a microhardness tester. A
SEM micrograph of the cross-section of the coating–substrate interface with microhardness indentations is shown in Fig. 7Ža.. X-ray elemental distribution of Al, Si, and Ti corresponding to
Table 4 Work of adhesion for metalrceramic systems Ceramic
Metallic melt
Surface energy, g 1y Õ of liquid–vapor interface ŽmJrm2 .
Contact angle in degrees
Work of adhesion, Wa ŽmJrm2 .
TiC
Al Si Mg Fe Al Si Mg Fe
914 860 583 1900 915 860 Not available 1785
118 32 50 39 160 34 Not available 49
485 1720 960 3578 55 1575 Not available 2955
TiB 2
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Fig. 7. Ža. SEM micrograph of the coating, coating–substrate interface, and substrate regions with indentations and corresponding X-ray elemental map of Žb. Al, Žc. Si, and Žd. Ti.
Fig. 7Ža. are represented in Fig. 7Žb., Žc. and Žd., respectively. The relative difference in size of microhardness indentations explains the variation in hardness. Considerably high microhardness values were observed in the coating. The two indentations taken in the coating showed different values of microhardness ŽKnoop., with the top part of the coating Ž572 " 60. relatively harder than the bottom part Ž454 " 30.. This is due to the presence of SiC, which apparently segregated to the top of the coating during laser processing. SiC, being less denser Ž3.1 to 3.2 grcm3 . in comparison to TiC Ž4.92 grcm3 ., floats to the top of molten coating during laser processing and remains at the top of the coating after resolidification. This fact is also evident from the corresponding X-ray elemental map of Si shown in Fig. 7Žc.. The hardness of the LMZ is 72 " 5 and that of unreacted substrate is 84 " 10. The relatively
low hardness value of LMZ is attributed to the dissolutionrredistribution of precipitates formed during the precipitation hardening of 6061 Al alloy. In the present study, attempts were made to semiquantitatively determine the interfacial strength. In the process, microhardness indentations were made
Table 5 SEM analysis of the indentations at the interface Applied load Žg.
Length of the longer diagonal Žmm.
SEM predictions
50 100 200 300 500 1000
32 61 72 107 116 150
No interfacial cracks No interfacial cracks No interfacial cracks No interfacial cracks No interfacial cracks No interfacial cracks
74
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along the interface at various loads Ž50, 100, 200, 300, 500, and 1000 g. to check for initiation of any cracks or delaminations at the site of indentations. The loads at which the microhardness tests are conducted and the corresponding SEM observations are illustrated in Table 5. Fig. 8 shows the corresponding
SEM micrographs of the indentations taken at all the test loads within the limitations of the indentor being used in the present study. SEM observations showed that there were no visible cracks or delaminations at these indentations. However, the possibility of fine stationary cracks within the coating is not ruled out
Fig. 8. SEM micrographs of the indentations taken at the coating–substrate interface at various loads.
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since this can only be observed at very high magnification. The dynamic behavior of these stationary cracks is being further investigated using a four-point bending test. Such tests will characterize the strength of ceramic–metal interface in a quantitative manner. 3.4. Tribological characterization 3.4.1. Wear tests The tribological performance of a ceramic coating depends on the characteristics such as hardness, thickness, internal stress level, and load-bearing capacity. When a coating cools from the high temperature of the melt zone in the deposition process, the internal compressive stress developed within the coating reduces the extent of surface fracture and, thus, increases the sliding wear resistance of the coated surface w29x. In addition, wear behavior of any material system is expected to be influenced by environment, such as dry Žair or inert atmosphere. and wet Žlubricated.. The thermo–physical Žforce and temperature. conditions at the point of contact are extreme during wear testing. These extreme conditions are favorable to cause degradationrdissociation Žphysical and chemical. of the material and, hence, occurrence of localized reactionŽs. between the material and surrounding environment, thereby providing different wear responses. Wear results in the loss of material, elevation of surface temperature, surface activation, and the removal of protective film, and, thus, may accelerate oxidation or corrosion w30x. Hence, it is important to
Fig. 9. Weight loss vs. time for the wear tests.
75
study the wear behavior of the coated samples to determine the nature of the coating. Fig. 9 shows a plot of weight loss vs. time for the wear tests conducted on TiC coated and uncoated Žsubstrate. surfaces of the laser-treated samples. It is clearly observed that the weight loss for the coated side was lower than the substrate side. The wear rate is also significantly improved. The total duration of the wear test was 20 min. However, while conducting the tests on the uncoated substrate side, the substantial weight losses were calculated for a time duration of only 30 s due to high seizure of the substrate material during sliding. Comparison of XRD analyses of the worn and unworn-coated surfaces as illustrated in Fig. 1 indicated that there are no new phases formed during the wear process. This suggested that the coating was chemically and physically stable under the dry sliding conditions used in the present study. Surface roughness profiles shown in Fig. 10Ža., Žb. and Žc. indicated that the average surface roughness, R a , of the test sample increased from 0.43 to 13.53 mm after laser processing. The average surface roughness of the worn surface is 6.51 mm. Topographic features of the worn surface were analyzed using SEM. Fig. 11Ža. shows a low magnification micrograph of the TiC coated 6061 Al sample. X-ray elemental distributions of Al, Ti, and Fe corresponding to Fig. 11Ža. are shown in Fig. 11Žb., Žc. and Žd., respectively. The worn surface of the coating is identified by continuous wear scars parallel to the laser tracks in some isolated regions of contact between the sample and wear-disc. During sliding, there is a material-to-material contact, followed by welding or fusing of the contacting asperities. When sliding continues, the asperities of the weaker material may shear off and transfer to the opposite surface w31x. In the present work, the hardened steel sliding disc Ž62 Rc., whose main constituent is Fe, shears off the Al present as the matrix of the coating, and Al is collected as the loose debris at the end of the scars. On the contrary, TiC present in the coating, being harder than Fe, shears off the material from the sliding disc, and, as a result, Fe is deposited in those scars. Such composite nature of wear and segregation of Al and Fe in particular regions of the worn surface are clearly seen from Fig. 11Ža., Žb. and Žd., respectively.
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Fig. 10. Optical profilometer figures showing the surface roughness of Ža. 6061 Al substrate, Žb. coated side and Žc. worn coated surface.
3.4.2. Coefficient of friction measurement The measurement of the coefficient of friction provides direct information about the work done to deform the surface of the material. However, it is not necessarily a direct indication of material loss or
separation of loose debris from the surface w15x. In the present study, the tribometer interfaced with a computer recorded the wear test parameters such as Voltage Ž V . and Current Ž I . as a function of time. Based on the principle of energy conservation, the
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77
Fig. 11. Ža. SEM micrograph of the worn surface and corresponding X-ray elemental map of Žb. Al, Žc. Ti and Žd. Fe.
frictional force is equal to the electrical work done by the motor given by
The frictional theory states that Wf s m NÕ
Wf s Voltage Ž DV . = Current Ž D I .
where N is the normal load, Õ is the linear speed of the sliding disc and m is the coefficient of friction. By equating the above equations, the coefficient of friction can be computed. The computed coefficient of friction for TiC coating has been plotted for the entire test time of 20 min as shown in Fig. 12. The best-fit line shows the coefficient of friction of the TiC-coated surface to be approximately 0.64.
Ž 8.
Ž 9.
4. Conclusions
Fig. 12. Coefficient of friction vs. time for the wear tests.
1. TiC has been deposited on 6061 Al alloy using the LSE technique. The coating is uniform, adherent, and free of cracks and porosities.
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2. The coating is composite in nature with the microstructure consisting primarily of hard TiC particles reinforced in a matrix of Al–Ti mixture. 3. The average surface roughness of the coating was found to be 13.5 mm. 4. Considerably high hardness values are obtained in the coating. 5. The coating is wear resistant and the wear resistance is highly influenced by the composite nature of the coating. 6. The coefficient of friction was computed to be approximately 0.64.
Acknowledgements Authors acknowledge the partial support from the Surface Technology Division of ALCOA.
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