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Fe nano-composite coating by laser surface engineering
Optics & Laser Technology 45 (2013) 647–653
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Optics & Laser Technology journal homepage: www.elsevier.com/locate/optlastec
Synthesis of TiB2–TiC/Fe nano-composite coating by laser surface engineering Baoshuai Du a,n, Sameer R. Paital b, Narendra B. Dahotre b a b
Shandong Electric Power Research Institute, Jinan 250002, China Laboratory of Laser Material Processing and Synthesis, Department of Materials Science and Engineering, University of North Texas, Denton, TX 76207, USA
a r t i c l e i n f o
abstract
Article history: Received 4 April 2012 Received in revised form 10 May 2012 Accepted 17 May 2012 Available online 5 June 2012
The study explores the synthesis of TiB2 and TiC reinforced Fe-based nano-composite coating by laser surface engineering using Ti, B4C and Fe powder mixture as precursor. The effect of laser scanning speed on the size, morphology, and amount of nano-sized ceramic reinforcements were studied at laser fluence of 1111 J/cm2, 1667 J/cm2 and 3333 J/cm2 respectively. A bimodal microstructure with TiC and TiB2 particles dispersed in fine a-Fe matrix was evolved in the laser processed coatings. Besides, the nature of formation of nano-sized ceramic phase was examined. The laser synthesized nano-composite coating yielded 3–5 times increase in microhardness. It appears that the presence of nano-sized TiB2 and TiC particles coupled with the highly refined a-Fe matrix improves the hardness significantly. This coating offers the potential to increase the hardness and toughness simultaneously for developing wear-resistance coatings. & 2012 Elsevier Ltd. All rights reserved.
Keywords: Laser surface engineering TiB2 TiC
1. Introduction Over the past several years various coating systems have been designed, fabricated and tested for use in wear-resistance applications [1–3]. Among these coatings, nano-composite coating with structure of hard ceramic phases embedded in a ductile matrix is an attractive candidate due to their high hardness, thermodynamic stability, chemical inertness, and high fracture toughness [4,5]. Laser surface engineering (LSE) which employs laser as the heat source to modify the surface of components has been proved to be a promising coating technique for fabricating wear-resistant coatings [6–10]. Moreover, the synthesis of in situ metal matrix composite (MMC) coatings using LSE has sparked a flurry of scientific interest [11–13]. This technique is highly efficient and flexible enough to produce in situ MMC coatings with desired reinforcements by employing different alloy systems. However, the dilemma is the presence of agglomeration and/or over-growth of ceramic particles when a high volume fraction of reinforcement is required, which leads to decrease in toughness of the coating. This may also result in cracking owing to the presence of residual stress generated during the solidification of the melt pool. Therefore, nano-composite coating that is characterized by high toughness seems promising in this sense.
n
Corresponding author. Tel.: þ86 531 68612562; fax: 86 531 68612560. E-mail address: [email protected] (B. Du).
0030-3992/$ - see front matter & 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.optlastec.2012.05.017
With development of advanced laser processing equipment and the understanding of coating design, the problem of agglomeration and over-growth of reinforcements can be circumvented to some extent by tailoring the laser processing parameters and the composition of precursor. Choudhury et al. synthesized nanostructured boride metal composite coatings by combining laser and Sol–gel technology [14]. They reveal that the coating consists of micro-level FeB phase and nano-particulate TiB phases in a (FeþFe2B) matrix. Nayak and Dahotre fabricated iron-oxide nanocoatings on aluminum cast alloys that is characterized by dispersion of Fe2O3 nano-particles in highly refined Al dendrite matrix by combining combustion synthesis and laser surface engineering [15]. In our previous work in the field of LSE [12], TiB2–TiC/Ni composite coating has been synthesized and it is found that TiB2 and TiC can be formed in situ through the reaction of Ti and B4C. The idea of synthesizing dual-ceramic phase reinforced composite coating is extended further by substituting nickel with iron in the present study, because the intended substitution of ecologically harmful and expensive cladding materials on cobalt and nickel by iron has attracted significant research interest for developing novel iron-based composite coatings for various applications [16,17]. The fabrication of TiB2–TiC/Fe nano-composite coating on steel substrate was accomplished by direct melting technique using LSE by introducing Ti, B4C and Fe mixture as precursor powder. One of the major advantages associated with the laser direct melting technique is control over the laser processing parameters to precisely mitigate laser material interaction and
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thereby synthesize a surface of uniform composition and morphology with highly refined microstructure [18]. Furthermore, compared with laser cladding or pulsed laser deposition (PLD), LSE is more versatile and robust because it neither needs the powder feeding process of laser cladding nor the vacuum atmosphere of PLD. In this study, the effect of laser fluence on the size, morphology, and amount of reinforcements were studied. The mechanism of formation of nano-sized reinforcements was also discussed. It is expected that this work will help to explore the synthesis and formation mechanism of many other nano-composite coatings using laser surface engineering technique.
Lens Laser beam
Precusor
v×t
2. Experimental Plates of AISI 1010 steel (0.08–0.13 wt% C, 0.3–0.6 wt% Mn, o0.04 wt% S, o0.05 wt% P, balance Fe) with dimensions of 75 mm 50 mm 6 mm were used as substrate during laser processing. They were polished using emery paper and then washed with acetone to provide a clean surface. Powder mixture of Fe (99.9% purity, 10 mm), titanium (99.5% purity, 10 mm) and B4C (99% purity, 10 mm) was used as the precursor with composition of 50Fe–36Ti–14B4C (wt%). In the Ti–B4C–Fe system, the relevant reaction of forming ceramic reinforcements is [19] 3Ti þ B4 C ¼ 2TiB2 þTiC
DG ¼ 740816 þ47:84T J=mol
E S
ð2Þ
where LF is the laser fluence, S is the area irradiated by laser beam, and E is the energy absorbed by the area of S, as shown in Fig. 1. Furthermore, E and S can be calculated as follows: E¼Pt
ð3Þ
S¼vtd
ð4Þ
where p, t and v are laser power, time throughout the area of S, and laser scanning speed respectively. By combining Eqs. 2–4, laser fluence can be calculated as following: LF ¼
P vd
S
Laser irradiated area
Fig. 1. Schematic of laser processing and the irradiated area.
ð1Þ
The Gibbs free energy change indicates that TiB2 and TiC have high tendency of formation in this Fe–Ti–B–C system, which has been proved in the literature [20]. Furthermore, the addition of Fe in the precursor can not only result in better coating appearance but also serve as the catalyst, because the eutectic reaction between Fe and Ti can lower the melting point of the precursor and facilitate the in situ reaction. The weight ratio of Ti to B4C in the precursor was chosen based on the above-mentioned reaction. Laser power and beam diameter used in the present study are 1000 W and 1 mm respectively. The details of the laser processing can be found in Ref. [12]. The selection of the beam diameter and laser power is based on the optimization of laser fluence which results in sufficient melting of the precursor and a portion of the substrate while the precursor material was not ablated by severe vaporization. It is well documented that the laser scanning speed determines the interaction time of laser beam with material and subsequently influences cooling rate and dilution ratio of the melt pool. Thus, during the experiment laser power was kept constant at 1000 W while laser scanning speed was varied at values of 30 mm/s, 60 mm/s and 90 mm/s, corresponding to laser fluence of 3333 J/cm2, 1667 J/cm2 and 1111 J/cm2 respectively, to study its effect on the structure and properties of coatings. Laser fluence is calculated by using the following equation: LF ¼
d
ð5Þ
Fig. 2. XRD spectra of the laser synthesized nano-composite coatings.
Structural and microstructure characterization were conducted by XRD and SEM respectively, the details of which can be found somewhere else [12]. In order to reveal the detailed microstructure, the nano-composite coating was also characterized by an H-800 transmission electron microscopy (TEM) at accelerating voltage of 150 kV. Slice sample for TEM analysis was cut parallel to the plane of the coating surface by low speed diamond saw. It was then polished by emery paper and finally thinning to perforation using a twin-jet polisher. Vickers indentation methods were employed to determine the hardness of the composite coating with a diamond Vickers tip, a 300 g load, and a dwell time of 12 s.
3. Results and Discussion 3.1. Phase analysis Results from the XRD characterization of the laser processed coatings are presented in Fig. 2. It clearly reveals that a dramatic phase evolution occurred after the precursor being irradiated by
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Fig. 4. Melt depth of the composite coating as a function of laser scanning speed. Fig. 3. Relative phase amount as a function of laser scanning speed.
the laser beam. TiB2, TiC, Fe3(B,C) and a-Fe were evolved in the coating. Thus, in comparing with the precursor, an inference can be made that TiB2 and TiC ceramic reinforcements were in situ formed under all the laser processing conditions employed in the present work. A more in-depth study was performed to identify semiquantitatively the amount of each phase presented in the coating. Relative amount of each phase was calculated by the following equation [21]: %Ii ¼
Ii
S31 Ii
ð6Þ
P where Ii is the phase in concern and 31 Ii corresponds to the sum of integrated peak intensity of each of the TiB2, TiC, a-Fe phases in the coating. It should be pointed out that Fe3(B,C) was excluded from the calculation due to its trace amount in the composite coating. The corresponding plans of peaks used in the calculation are labeled in Fig. 2. Moreover, it should be noted that the value obtained here can only represent the relative amount of each phase instead of the actual proportion of it. The relative amounts of these phases as function of laser scan speed are presented in Fig. 3 and indicate that the amount of TiB2 and TiC increase with the increase of laser scanning speed while a-Fe shows an opposite trend. While the relative phase amount of TiB2 increases from 15.9% to 24.2% and that of TiC from 7.1% to 8.6% for samples processed with scanning speed of 30 mm/s and 90 mm/s respectively, relative amount of a-Fe declines from 77% to 67.2% for the same samples. The variation of amount of relative phase can be understood by considering the dilution effect caused by the melting of substrate during the laser surface treatment. During laser processing a portion of the substrate is melted and mixed with the precursor to form a melting pool which subsequently solidifies and results in the final coating. Thus, the ratio of molten substrate to the precursor (dilution ratio) determines the relative amount of phases presented in the coating. The calculation procedure of dilution ratio can be found in Ref. [12]. Dilution ratios are 79%, 72%, and 66% for samples processed with 30 mm/s, 60 mm/s, and 90 mm/s respectively. For the current study, laser scanning speed was varied during the laser surface treatment. High laser scan speed leads to short laser-material interaction time (laser beam dwell time) and less energy input per unit volume. Therefore, less melt depth was obtained for samples processed with higher scanning speed, as
indicated in Fig. 4. This also means less substrate is melted and mixed in the melting pool, in other words dilution ratio is lower for samples processed with higher scanning speed. As a result, volume fraction of TiB2 and TiC in the composite coating is higher for sample processed with higher laser scanning speed (less laser fluence), while that of a-Fe shows an opposite trend. 3.2. Microstructure evolution Fig. 5(a)–(h) illustrate the microstructure of laser synthesized coating. The low magnification view of the laser processed coating shows that the coating is metallurgically-bonded, adherent, continuous, and pore and crack-free (Fig. 5(a)). The low carbon steel region adjacent to the fusion line exhibits a martensitic structure, because this region experiences rapid cooling due to the heat sink effect of the substrate during the laser treatment. Furthermore, it is obvious that the coating is composite in nature, which shows randomly distributed particles embedded in the fine-grained matrix (Fig. 5(b)). Combining with the XRD results, it is obvious that the reinforcements are TiB2 and TiC, while the fine-grained matrix is a-Fe. According to the image analysis conducted using a public-domain software, ImageJ, size of the reinforcements observed in the composite coating has an average diameter of 171 nm, 175 nm, and 119 nm for samples processed with scanning speed of 30 mm/s, 60 mm/s, and 90 mm/s respectively. Occasionally, particles with size less than 50 nm were also detected in the composite coating. Furthermore, it can be found that the matrix of a-Fe is also highly refined (less than 1 mm in average grain size). Since the samples processed under varying laser scan speeds showed no significant change in the nature of microstructure, the sample processed with 60 mm/s was selected for TEM observations (Fig. 6). The low magnification TEM image (Fig. 6(a)) clearly indicates dispersion of white nano-particles in the gray and dark matrix due to the diffraction contrast. Dislocation lines and network induced by the rapid cooling process can also be detected, as shown by the white arrow in Fig. 6(b). Clean and adherent interface between the nano-sized particles and a-Fe can be distinguished, which is attributed to the fact that these particles were formed by the in situ reaction. Moreover, as indicated by the white polygon in Fig. 6(c) and (d), typical reinforcements exhibits hexagonal and rhomboidal shape corresponding to TiB2 and TiC crystal structures respectively, which will be discussed later. Understanding the formation mechanism of ceramic reinforcements plays a key role in tailoring and optimization of the
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Melt zone
Fusing line Heat affected zone
100 µm
10 µm
1 µm
500nm
1 µm
500nm
1 µm
200nm
Fig. 5. SEM images of the cross-sectional view of laser synthesized nano-composite coatings: (a) low magnification view of the melt profile (60 mm/s); (b) low magnification view of the melt zone (60 mm/s) and inset showing the schematic of the composite microstructure; (c) and (d) high magnification views of coating (30 mm/s) sample processed with 30 mm/s; (e) and (f) high magnification views of coating (60 mm/s); (g) and (h) high magnification views of coating (90 mm/s).
structure and properties of the nano-composite coating. During laser surface engineering, laser beam with high irradiance (fluence) imparts large amount of thermal energy into the surface of the sample and causes the precursor (Fe, B4C and Ti) and a thin layer near the surface of the steel substrate to melt and mix together. Furthermore, the exothermic reaction between B4C and Ti can (Eq. (1), DH0 ¼ 686 kJ) contribute to the heating of the material. As a result, a confined volume of the surface region forms a melt pool. Since the rest of massive substrate volume remains cool and acts as a heat sink, a very large temperature gradient is established and high cooling rate is induced in the melt pool. During the subsequent cooling, the high melting reinforcements of TiB2 and TiC precipitate from the liquid melt pool first and grow into the final shape. The remaining liquid solidifies consequently and transforms into a-Fe and Fe3(B, C) after the formation of TiB2 and TiC. The high cooling rate causes remarkable undercooling that in turn induces
high nucleation rate for the reinforcements. Therefore, refinement of TiB2 and TiC particles and matrix grain can be achieved for laser processed coating. As stated earlier, sample processed with 90 mm/s possessed the finest particle size and samples processed with 30 mm/s and 60 mm/s possessed the particles with marginal average size difference. This can be understood by taking into consideration both the cooling rate and dilution ratio of the coating. It is obvious that higher scan speed is beneficial for obtaining a high cooling rate in the melt pool. Thus, the resultant large undercooling induces higher nucleation rate and less time for the particles to grow for sample processed with 90 mm/s. On the contrary, the relative lower cooling rate of samples processed with 60 mm/s and 30 mm/s make the in situ formed particles of these samples larger than that of 90 mm/s. However, although samples processed with 60 mm/s cool at higher cooling rates than that of 30 mm/s, the dilution ratios in this sample is
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Dislocation
110
2µm
200
1µm
110
200nm
200nm
Fig. 6. TEM images of the sample processed with 60 mm/s: (a) and (b) low magnification view of the nano-composite coating and inset in (b) is the SAD pattern of a-Fe matrix ([0 0 1] axis); (c) and (d) high magnification view of the nano-composite coating showing the in situ synthesized TiB2 and TiC particles.
lower (79% and 72% for samples processed with 30 mm/s and 60 mm/s respectively) leading to retention of higher content of reinforcements-forming elements in the melt pool. As a result, the driving force for particle growth is higher and samples processed with 30 mm/s and 60 mm/s contained similar particle size. Thus, the competition between dilution ratio and cooling rate determines the final particle size. The embedding particles of nano-sized hard ceramic phase (TiB2 and TiC) into a soft metallic matrix allows for high ductility and thus leading to high toughness of the coating. It should be also noted that the grain size of the Fe matrix is also highly refined (scaling down to 1 mm). The grain-refining effect provides dislocation and crack splitting, resulting in improved ductility and toughness for the nano-composite coating. As the morphology of the reinforcements affect the final properties of the composite coating, it is desirable to understand the morphology evolution of TiB2 and TiC reinforcements. Based on standard argument, the thermodynamically preferred crystal morphology is defined by the minimization of surface free energy Gs, which is defined as following [21]: X Gs ¼ ghkl Ahkl ð7Þ hkl
where ghkl is the surface energy per unit area of a given polyhedral face, and Ahkl is the area of the face. TiB2 has a hexagonal AlB2type structure with a P6/mmm space group involving alternating planes of titanium and boron atoms (Fig. 7) [22]. Past calculations for attachment energies in TiB2 crystals reveal the following hierarchy of growth rates (GR) [23]: GR0001 oGR1010 o GR1011 oGR1210 o GR1211
ð8Þ
It has been theoretically confirmed that (0 0 0 1) plane is the close-packed plane and {1 0 1 0} are the less stable planes for TiB2. Thus, the hexagonal plates of TiB2 in the composite coating reflect its crystal habit. On the contrary, Bates et al. [22] observed size-
related morphologies for TiB2 due to different surface-to-volume ratio and surface energy per unit area, which shows that fine particles (5–10 nm) appeared to have equiaxed, cuboidal (or rohombic) morphologies while larger particles (more than 20 nm) were hexagonal plates [21], and their result is consistent with the observations of the present study. TiC has a rock-salt NaCl-type structure that Titanium atoms array the face-centered cubic (fcc) sublattice while carbon atoms occupy the free sites of the octahedral configuration forming another fcc sublattice (Fig. 7). {1 1 1} planes of TiC share the highest surface atomic density and the lowest surface energy. Thus, equilibrium TiC crystal shows octahedral shape enclosed by {1 1 1} planes. Besides, truncated octahedral shape of TiC can be also formed due to the difference of growth rate in the o0 0 14 and o1 1 14 directions. Thus, TiC particles tend to show octahedron and/or truncated octahedral shape. The rhomboid and cubic shape of TiC shown in our study is due to the fact that the surface of sample for observation intersects with the TiC crystal with a certain angle.
3.3. Microhardness By taking measurement using Vikers indenter, microhardness was obtained for the laser processed coatings. It demonstrates that high hardness can be imparted to the nano-composite coating (Fig. 8). 3–5 times increase of hardness can be observed for the nano-composite coatings compared with that of the steel substrate. The average hardness values of the coatings are 498 kg/ mm2, 596 kg/mm2, and 689 kg/mm2 for samples processed with 30 mm/s, 60 mm/s, and 90 mm/s respectively. In general, the increase of hardness with increasing laser scanning speed is attributed to the higher content of reinforcements presented in the coating with higher laser processing speed, as indicated by the XRD analysis. Besides the presence of nano-sized TiB2 and TiC
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Fig. 7. (a) Schematic of crystal structure of TiB2; (b) projection of Ti and B atoms along the [0 0 0 1] direction for TiB2 crystal; (c) schematic of crystal structure of TiC; and (d) projection of Ti and C atoms along the [0 0 1] direction for TiC crystal.
improvement of hardness of the nano-composite coating will lead to the increase of wear resistance. Moreover, the in situ formed particles promote strong particle–matrix bonding, thereby reducing pull-out of particles during wear. High toughness is also desirable for materials designed for wear application since it can prevent crack-forming and spallation. It is obvious that the rapid cooling rate associated with laser processing induces the microstructural refinement for both the reinforcements and matrix grains, and reduction of the size of the reinforcements to nanometer size offers the unique opportunity to design wear resistant coating with optimal combination of hardness and toughness [25]. Nevertheless, a comparison of the present study with Ref. [12] shows the overall hardness of TiB2–TiC/Fe is lower than that of TiB2–TiC/Ni. It is attributed to the low volume fraction of reinforcements in the TiB2–TiC/Fe composite coating compared with the TiB2–TiC/Ni, which may result from the complex laser-materials interaction. On the other hand, nickel element plays the role of solid solution to increase the hardness of the matrix in the composite coating. This also gives indication that by introducing alloy elements to the iron matrix through the solid–solution strengthening mechanism the hardness of the composite coating can be further increased.
Fig. 8. Microhardness profiles of the laser processed samples.
reinforcement, the grain refinement effect induced by the a-Fe can also contribute to the increase of hardness. The wear resistance can often be related to hardness through the Archard equation, which states the volume of material worn is inversely proportional to hardness for a given set of test conditions [24]: V ¼K l
P H
ð9Þ
where K is the wear coefficient, l is the sliding distance, P is the applied load and H is the hardness. Thus, it is apparent that the
4. Conclusion TiB2 and TiC reinforced Fe-based nano-composite coating was in situ fabricated by laser surface engineering using Ti, B4C and Fe as the components of precursor. Structural characterization reveals that the composite coating consists of TiB2, TiC, Fe3(B,C) and a-Fe. By comparing integrated intensity of each phase, it is found that the relative amount of TiB2 and TiC increase with the increasing laser scanning speed which was attributed to the
B. Du et al. / Optics & Laser Technology 45 (2013) 647–653
variation of dilution ratio. Microstructure of the coating indicates that nano-sized TiB2 and TiC particles uniformly distribute in the fine-grained a-Fe matrix. The dual-ceramic phase design is effective in reducing the size of in situ formed reinforcements. Reinforcements of TiB2 and TiC formed following the liquidprecipitation mechanism and their morphology is determined by the crystallographic habit and minimization of surface free energy. Owing to the presence of nano-sized reinforcements (TiB2 and TiC) and the highly refined a-Fe matrix compared with the steel substrate, microhardness value of the composite coating improved significantly (498 kg/mm2, 596 kg/mm2, and 689 kg/mm2 for samples processed with 30 mm/s, 60 mm/s, and 90 mm/s respectively). The preliminary evaluation based on microstructural and hardness analyses show that nano TiB2 and TiC reinforced iron coating using LSE appear to hold promise in wear application. Future work would involve alloying the iron-matrix of the composite coating to further improve its hardness and evaluation of the wear resistance of the nano-composite coating. References [1] Degnan CC, Shipway PH. A comparison of the reciprocating sliding wear behaviour of steel based metal matrix composites processed from selfpropagating high-temperature synthesized Fe–TiC and Fe–TiB2 master alloys. Wear 2002;252:832–41. [2] Zhou Z, Ross IM, Ma L, Rainforth WM, Ehiasarian AP, Hovsepian P. Wear of hydrogen free C/Cr PVD coating against Al2O3 at room temperature. Wear 2011;271:2150–6. [3] Hadad M, Hitzek R, Buergler P, Rohr L, Siegmann S. Wear performance of sandwich structured WC–Co–Cr thermally sprayed coatings using different intermediate layers. Wear 2007;263:691–9. [4] Voevodin AA, Zabinski JS. Nanocomposite and nanostructured tribological materials for space applications. Composites Science and Technology 2005;65: 741–9. [5] Voevodin AA, Hu JJ, Fitz TA, Zabinski JS. Tribological properties of adaptive nanocomposite coatings made of yttria stabilized zirconia and gold. Surface and Coatings Technology 2001;146–147:351–6. [6] Harimkar SP, Dahotre NB. Rapid surface microstructuring of porous alumina ceramic using continuous wave Nd:YAG laser. Journal of Materials Processing Technology 2009;209:4744–9. [7] Harimkar SP, Dahotre NB. Characterization of microstructure in laser surface modified alumina ceramic. Materials Characterization 2008;59:700–7. [8] Agarwal A, Dahotre NB. Comparative wear in titanium diboride coatings on steel using high energy density processes. Wear 2000;240:144–51. [9] Samant AN, Paital SR, Dahotre NB. Process optimization in laser surface structuring of alumina. Journal of Materials Processing Technology 2008;203:498–504.
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[10] Katakam S, Hwang JY, Vora H, Harimkar SP, Banerjee R, Dahotre NB. Laserinduced thermal and spatial nanocrystallization of amorphous Fe–Si–B alloy. Scripta Materialia 2012;66:538–41. [11] Gopagoni S, Hwang JY, Singh ARP, Mensah BA, Bunce N, Tiley J, Scharf TW, Banerjee R. Microstructural evolution in laser deposited nickel-titaniumcarbon in situ metal matrix composites. Journal of Alloys and Compounds 2011;509:1255–60. [12] Du B, Paital SR, Dahotre NB. Phase constituents and microstructure of laser synthesized TiB2–TiC reinforced composite coating on steel. Scripta Materialia 2008;59:1147–50. [13] Man HC, Zhang S, Cheng FT, Yue TM. Microstructure and formation mechanism of in situ synthesized TiC/Ti surface MMC on Ti–6Al–4V by laser cladding. Scripta Materialia 2001;44:2801–7. [14] Choudhury AR, Ezz T, Li L. Synthesis of hard nano-structured metal matrix composite boride coatings using combined laser and sol–gel technology. Materials Science and Engineering: A 2007;445–446:193–202. [15] Nayak S, Dahotre NB. The laser-induced combustion synthesis of iron-oxide nanocomposite coatings on aluminum. Journal of the Minerals, Metals and Materials Society 2002;54:30–41. [16] Lampke T, Wielage B, Pokhmurska H, Rupprecht C, Schberth S, Drehmann R, Schreiber F. Development of particle-reinforced nanostructured iron-based composite alloys for thermal spraying. Surface and Coatings Technology 2011;205:3671–6. [17] Du B, Wang X, Zou Z. Microstructure and tribological behavior of laser in situ synthesized TiC-reinforced Fe-based composite coatings. Tribology Letters 2011;43:295–301. [18] Paital SR, Bhattacharya A, Moncayo M, Ho YH, Mahdak K, Nag S, Banerjee R, Dahotre NB. Improved corrosion and wear resistance of Mg alloys via laser surface modification of Al on AZ31B. Surface and Coatings Technology 2012;206:2308–15. [19] Liang Y, Che Y. Notebook of thermodynamic data of inorganic. Shenyang: East-north University Press; 1996. [20] Wang Z, Zhou X, Zhao G. Microstructure and formation mechanism of in-situ TiC–TiB2/Fe composite coating. Transactions of Nonferrous Metals Society of China 2008;18:831–5. [21] Greish Y, Brown P. Phase evolution during the formation of stoichiometric hydroxyapatite at 37.4 1C. Journal of Biomedical Materials Research 2003;67:632–7. [22] Bates SE, Buhro WE, Frey CA, Sastry. SML, Kelton KF. Synthesis of titanium boride TiB2 nanocrystallites by solution-phase processing. Journal of Materials Research 1995;10:2599–612. [23] Vallauri D, Adria´n ICA, Chrysanthou A. TiC–TiB2 composites: A review of phase relationships, processing and properties. Journal of the European Ceramic Society 2009;28:1697–713. [24] Rupert TJ, Schuh CA. Sliding wear of nanocrystalline Ni–W: structural evolution and the apparent breakdown of Archard scaling. Acta Materialia 2010;58:4137–48. [25] Tjong SC, Ma ZY. Microstructural and mechanical characteristics of in situ metal matrix composites. Materials Science and Engineering R 2000;29: 49–113.
Contents lists available at SciVerse ScienceDirect
Optics & Laser Technology journal homepage: www.elsevier.com/locate/optlastec
Synthesis of TiB2–TiC/Fe nano-composite coating by laser surface engineering Baoshuai Du a,n, Sameer R. Paital b, Narendra B. Dahotre b a b
Shandong Electric Power Research Institute, Jinan 250002, China Laboratory of Laser Material Processing and Synthesis, Department of Materials Science and Engineering, University of North Texas, Denton, TX 76207, USA
a r t i c l e i n f o
abstract
Article history: Received 4 April 2012 Received in revised form 10 May 2012 Accepted 17 May 2012 Available online 5 June 2012
The study explores the synthesis of TiB2 and TiC reinforced Fe-based nano-composite coating by laser surface engineering using Ti, B4C and Fe powder mixture as precursor. The effect of laser scanning speed on the size, morphology, and amount of nano-sized ceramic reinforcements were studied at laser fluence of 1111 J/cm2, 1667 J/cm2 and 3333 J/cm2 respectively. A bimodal microstructure with TiC and TiB2 particles dispersed in fine a-Fe matrix was evolved in the laser processed coatings. Besides, the nature of formation of nano-sized ceramic phase was examined. The laser synthesized nano-composite coating yielded 3–5 times increase in microhardness. It appears that the presence of nano-sized TiB2 and TiC particles coupled with the highly refined a-Fe matrix improves the hardness significantly. This coating offers the potential to increase the hardness and toughness simultaneously for developing wear-resistance coatings. & 2012 Elsevier Ltd. All rights reserved.
Keywords: Laser surface engineering TiB2 TiC
1. Introduction Over the past several years various coating systems have been designed, fabricated and tested for use in wear-resistance applications [1–3]. Among these coatings, nano-composite coating with structure of hard ceramic phases embedded in a ductile matrix is an attractive candidate due to their high hardness, thermodynamic stability, chemical inertness, and high fracture toughness [4,5]. Laser surface engineering (LSE) which employs laser as the heat source to modify the surface of components has been proved to be a promising coating technique for fabricating wear-resistant coatings [6–10]. Moreover, the synthesis of in situ metal matrix composite (MMC) coatings using LSE has sparked a flurry of scientific interest [11–13]. This technique is highly efficient and flexible enough to produce in situ MMC coatings with desired reinforcements by employing different alloy systems. However, the dilemma is the presence of agglomeration and/or over-growth of ceramic particles when a high volume fraction of reinforcement is required, which leads to decrease in toughness of the coating. This may also result in cracking owing to the presence of residual stress generated during the solidification of the melt pool. Therefore, nano-composite coating that is characterized by high toughness seems promising in this sense.
n
Corresponding author. Tel.: þ86 531 68612562; fax: 86 531 68612560. E-mail address: [email protected] (B. Du).
0030-3992/$ - see front matter & 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.optlastec.2012.05.017
With development of advanced laser processing equipment and the understanding of coating design, the problem of agglomeration and over-growth of reinforcements can be circumvented to some extent by tailoring the laser processing parameters and the composition of precursor. Choudhury et al. synthesized nanostructured boride metal composite coatings by combining laser and Sol–gel technology [14]. They reveal that the coating consists of micro-level FeB phase and nano-particulate TiB phases in a (FeþFe2B) matrix. Nayak and Dahotre fabricated iron-oxide nanocoatings on aluminum cast alloys that is characterized by dispersion of Fe2O3 nano-particles in highly refined Al dendrite matrix by combining combustion synthesis and laser surface engineering [15]. In our previous work in the field of LSE [12], TiB2–TiC/Ni composite coating has been synthesized and it is found that TiB2 and TiC can be formed in situ through the reaction of Ti and B4C. The idea of synthesizing dual-ceramic phase reinforced composite coating is extended further by substituting nickel with iron in the present study, because the intended substitution of ecologically harmful and expensive cladding materials on cobalt and nickel by iron has attracted significant research interest for developing novel iron-based composite coatings for various applications [16,17]. The fabrication of TiB2–TiC/Fe nano-composite coating on steel substrate was accomplished by direct melting technique using LSE by introducing Ti, B4C and Fe mixture as precursor powder. One of the major advantages associated with the laser direct melting technique is control over the laser processing parameters to precisely mitigate laser material interaction and
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thereby synthesize a surface of uniform composition and morphology with highly refined microstructure [18]. Furthermore, compared with laser cladding or pulsed laser deposition (PLD), LSE is more versatile and robust because it neither needs the powder feeding process of laser cladding nor the vacuum atmosphere of PLD. In this study, the effect of laser fluence on the size, morphology, and amount of reinforcements were studied. The mechanism of formation of nano-sized reinforcements was also discussed. It is expected that this work will help to explore the synthesis and formation mechanism of many other nano-composite coatings using laser surface engineering technique.
Lens Laser beam
Precusor
v×t
2. Experimental Plates of AISI 1010 steel (0.08–0.13 wt% C, 0.3–0.6 wt% Mn, o0.04 wt% S, o0.05 wt% P, balance Fe) with dimensions of 75 mm 50 mm 6 mm were used as substrate during laser processing. They were polished using emery paper and then washed with acetone to provide a clean surface. Powder mixture of Fe (99.9% purity, 10 mm), titanium (99.5% purity, 10 mm) and B4C (99% purity, 10 mm) was used as the precursor with composition of 50Fe–36Ti–14B4C (wt%). In the Ti–B4C–Fe system, the relevant reaction of forming ceramic reinforcements is [19] 3Ti þ B4 C ¼ 2TiB2 þTiC
DG ¼ 740816 þ47:84T J=mol
E S
ð2Þ
where LF is the laser fluence, S is the area irradiated by laser beam, and E is the energy absorbed by the area of S, as shown in Fig. 1. Furthermore, E and S can be calculated as follows: E¼Pt
ð3Þ
S¼vtd
ð4Þ
where p, t and v are laser power, time throughout the area of S, and laser scanning speed respectively. By combining Eqs. 2–4, laser fluence can be calculated as following: LF ¼
P vd
S
Laser irradiated area
Fig. 1. Schematic of laser processing and the irradiated area.
ð1Þ
The Gibbs free energy change indicates that TiB2 and TiC have high tendency of formation in this Fe–Ti–B–C system, which has been proved in the literature [20]. Furthermore, the addition of Fe in the precursor can not only result in better coating appearance but also serve as the catalyst, because the eutectic reaction between Fe and Ti can lower the melting point of the precursor and facilitate the in situ reaction. The weight ratio of Ti to B4C in the precursor was chosen based on the above-mentioned reaction. Laser power and beam diameter used in the present study are 1000 W and 1 mm respectively. The details of the laser processing can be found in Ref. [12]. The selection of the beam diameter and laser power is based on the optimization of laser fluence which results in sufficient melting of the precursor and a portion of the substrate while the precursor material was not ablated by severe vaporization. It is well documented that the laser scanning speed determines the interaction time of laser beam with material and subsequently influences cooling rate and dilution ratio of the melt pool. Thus, during the experiment laser power was kept constant at 1000 W while laser scanning speed was varied at values of 30 mm/s, 60 mm/s and 90 mm/s, corresponding to laser fluence of 3333 J/cm2, 1667 J/cm2 and 1111 J/cm2 respectively, to study its effect on the structure and properties of coatings. Laser fluence is calculated by using the following equation: LF ¼
d
ð5Þ
Fig. 2. XRD spectra of the laser synthesized nano-composite coatings.
Structural and microstructure characterization were conducted by XRD and SEM respectively, the details of which can be found somewhere else [12]. In order to reveal the detailed microstructure, the nano-composite coating was also characterized by an H-800 transmission electron microscopy (TEM) at accelerating voltage of 150 kV. Slice sample for TEM analysis was cut parallel to the plane of the coating surface by low speed diamond saw. It was then polished by emery paper and finally thinning to perforation using a twin-jet polisher. Vickers indentation methods were employed to determine the hardness of the composite coating with a diamond Vickers tip, a 300 g load, and a dwell time of 12 s.
3. Results and Discussion 3.1. Phase analysis Results from the XRD characterization of the laser processed coatings are presented in Fig. 2. It clearly reveals that a dramatic phase evolution occurred after the precursor being irradiated by
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Fig. 4. Melt depth of the composite coating as a function of laser scanning speed. Fig. 3. Relative phase amount as a function of laser scanning speed.
the laser beam. TiB2, TiC, Fe3(B,C) and a-Fe were evolved in the coating. Thus, in comparing with the precursor, an inference can be made that TiB2 and TiC ceramic reinforcements were in situ formed under all the laser processing conditions employed in the present work. A more in-depth study was performed to identify semiquantitatively the amount of each phase presented in the coating. Relative amount of each phase was calculated by the following equation [21]: %Ii ¼
Ii
S31 Ii
ð6Þ
P where Ii is the phase in concern and 31 Ii corresponds to the sum of integrated peak intensity of each of the TiB2, TiC, a-Fe phases in the coating. It should be pointed out that Fe3(B,C) was excluded from the calculation due to its trace amount in the composite coating. The corresponding plans of peaks used in the calculation are labeled in Fig. 2. Moreover, it should be noted that the value obtained here can only represent the relative amount of each phase instead of the actual proportion of it. The relative amounts of these phases as function of laser scan speed are presented in Fig. 3 and indicate that the amount of TiB2 and TiC increase with the increase of laser scanning speed while a-Fe shows an opposite trend. While the relative phase amount of TiB2 increases from 15.9% to 24.2% and that of TiC from 7.1% to 8.6% for samples processed with scanning speed of 30 mm/s and 90 mm/s respectively, relative amount of a-Fe declines from 77% to 67.2% for the same samples. The variation of amount of relative phase can be understood by considering the dilution effect caused by the melting of substrate during the laser surface treatment. During laser processing a portion of the substrate is melted and mixed with the precursor to form a melting pool which subsequently solidifies and results in the final coating. Thus, the ratio of molten substrate to the precursor (dilution ratio) determines the relative amount of phases presented in the coating. The calculation procedure of dilution ratio can be found in Ref. [12]. Dilution ratios are 79%, 72%, and 66% for samples processed with 30 mm/s, 60 mm/s, and 90 mm/s respectively. For the current study, laser scanning speed was varied during the laser surface treatment. High laser scan speed leads to short laser-material interaction time (laser beam dwell time) and less energy input per unit volume. Therefore, less melt depth was obtained for samples processed with higher scanning speed, as
indicated in Fig. 4. This also means less substrate is melted and mixed in the melting pool, in other words dilution ratio is lower for samples processed with higher scanning speed. As a result, volume fraction of TiB2 and TiC in the composite coating is higher for sample processed with higher laser scanning speed (less laser fluence), while that of a-Fe shows an opposite trend. 3.2. Microstructure evolution Fig. 5(a)–(h) illustrate the microstructure of laser synthesized coating. The low magnification view of the laser processed coating shows that the coating is metallurgically-bonded, adherent, continuous, and pore and crack-free (Fig. 5(a)). The low carbon steel region adjacent to the fusion line exhibits a martensitic structure, because this region experiences rapid cooling due to the heat sink effect of the substrate during the laser treatment. Furthermore, it is obvious that the coating is composite in nature, which shows randomly distributed particles embedded in the fine-grained matrix (Fig. 5(b)). Combining with the XRD results, it is obvious that the reinforcements are TiB2 and TiC, while the fine-grained matrix is a-Fe. According to the image analysis conducted using a public-domain software, ImageJ, size of the reinforcements observed in the composite coating has an average diameter of 171 nm, 175 nm, and 119 nm for samples processed with scanning speed of 30 mm/s, 60 mm/s, and 90 mm/s respectively. Occasionally, particles with size less than 50 nm were also detected in the composite coating. Furthermore, it can be found that the matrix of a-Fe is also highly refined (less than 1 mm in average grain size). Since the samples processed under varying laser scan speeds showed no significant change in the nature of microstructure, the sample processed with 60 mm/s was selected for TEM observations (Fig. 6). The low magnification TEM image (Fig. 6(a)) clearly indicates dispersion of white nano-particles in the gray and dark matrix due to the diffraction contrast. Dislocation lines and network induced by the rapid cooling process can also be detected, as shown by the white arrow in Fig. 6(b). Clean and adherent interface between the nano-sized particles and a-Fe can be distinguished, which is attributed to the fact that these particles were formed by the in situ reaction. Moreover, as indicated by the white polygon in Fig. 6(c) and (d), typical reinforcements exhibits hexagonal and rhomboidal shape corresponding to TiB2 and TiC crystal structures respectively, which will be discussed later. Understanding the formation mechanism of ceramic reinforcements plays a key role in tailoring and optimization of the
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Melt zone
Fusing line Heat affected zone
100 µm
10 µm
1 µm
500nm
1 µm
500nm
1 µm
200nm
Fig. 5. SEM images of the cross-sectional view of laser synthesized nano-composite coatings: (a) low magnification view of the melt profile (60 mm/s); (b) low magnification view of the melt zone (60 mm/s) and inset showing the schematic of the composite microstructure; (c) and (d) high magnification views of coating (30 mm/s) sample processed with 30 mm/s; (e) and (f) high magnification views of coating (60 mm/s); (g) and (h) high magnification views of coating (90 mm/s).
structure and properties of the nano-composite coating. During laser surface engineering, laser beam with high irradiance (fluence) imparts large amount of thermal energy into the surface of the sample and causes the precursor (Fe, B4C and Ti) and a thin layer near the surface of the steel substrate to melt and mix together. Furthermore, the exothermic reaction between B4C and Ti can (Eq. (1), DH0 ¼ 686 kJ) contribute to the heating of the material. As a result, a confined volume of the surface region forms a melt pool. Since the rest of massive substrate volume remains cool and acts as a heat sink, a very large temperature gradient is established and high cooling rate is induced in the melt pool. During the subsequent cooling, the high melting reinforcements of TiB2 and TiC precipitate from the liquid melt pool first and grow into the final shape. The remaining liquid solidifies consequently and transforms into a-Fe and Fe3(B, C) after the formation of TiB2 and TiC. The high cooling rate causes remarkable undercooling that in turn induces
high nucleation rate for the reinforcements. Therefore, refinement of TiB2 and TiC particles and matrix grain can be achieved for laser processed coating. As stated earlier, sample processed with 90 mm/s possessed the finest particle size and samples processed with 30 mm/s and 60 mm/s possessed the particles with marginal average size difference. This can be understood by taking into consideration both the cooling rate and dilution ratio of the coating. It is obvious that higher scan speed is beneficial for obtaining a high cooling rate in the melt pool. Thus, the resultant large undercooling induces higher nucleation rate and less time for the particles to grow for sample processed with 90 mm/s. On the contrary, the relative lower cooling rate of samples processed with 60 mm/s and 30 mm/s make the in situ formed particles of these samples larger than that of 90 mm/s. However, although samples processed with 60 mm/s cool at higher cooling rates than that of 30 mm/s, the dilution ratios in this sample is
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Dislocation
110
2µm
200
1µm
110
200nm
200nm
Fig. 6. TEM images of the sample processed with 60 mm/s: (a) and (b) low magnification view of the nano-composite coating and inset in (b) is the SAD pattern of a-Fe matrix ([0 0 1] axis); (c) and (d) high magnification view of the nano-composite coating showing the in situ synthesized TiB2 and TiC particles.
lower (79% and 72% for samples processed with 30 mm/s and 60 mm/s respectively) leading to retention of higher content of reinforcements-forming elements in the melt pool. As a result, the driving force for particle growth is higher and samples processed with 30 mm/s and 60 mm/s contained similar particle size. Thus, the competition between dilution ratio and cooling rate determines the final particle size. The embedding particles of nano-sized hard ceramic phase (TiB2 and TiC) into a soft metallic matrix allows for high ductility and thus leading to high toughness of the coating. It should be also noted that the grain size of the Fe matrix is also highly refined (scaling down to 1 mm). The grain-refining effect provides dislocation and crack splitting, resulting in improved ductility and toughness for the nano-composite coating. As the morphology of the reinforcements affect the final properties of the composite coating, it is desirable to understand the morphology evolution of TiB2 and TiC reinforcements. Based on standard argument, the thermodynamically preferred crystal morphology is defined by the minimization of surface free energy Gs, which is defined as following [21]: X Gs ¼ ghkl Ahkl ð7Þ hkl
where ghkl is the surface energy per unit area of a given polyhedral face, and Ahkl is the area of the face. TiB2 has a hexagonal AlB2type structure with a P6/mmm space group involving alternating planes of titanium and boron atoms (Fig. 7) [22]. Past calculations for attachment energies in TiB2 crystals reveal the following hierarchy of growth rates (GR) [23]: GR0001 oGR1010 o GR1011 oGR1210 o GR1211
ð8Þ
It has been theoretically confirmed that (0 0 0 1) plane is the close-packed plane and {1 0 1 0} are the less stable planes for TiB2. Thus, the hexagonal plates of TiB2 in the composite coating reflect its crystal habit. On the contrary, Bates et al. [22] observed size-
related morphologies for TiB2 due to different surface-to-volume ratio and surface energy per unit area, which shows that fine particles (5–10 nm) appeared to have equiaxed, cuboidal (or rohombic) morphologies while larger particles (more than 20 nm) were hexagonal plates [21], and their result is consistent with the observations of the present study. TiC has a rock-salt NaCl-type structure that Titanium atoms array the face-centered cubic (fcc) sublattice while carbon atoms occupy the free sites of the octahedral configuration forming another fcc sublattice (Fig. 7). {1 1 1} planes of TiC share the highest surface atomic density and the lowest surface energy. Thus, equilibrium TiC crystal shows octahedral shape enclosed by {1 1 1} planes. Besides, truncated octahedral shape of TiC can be also formed due to the difference of growth rate in the o0 0 14 and o1 1 14 directions. Thus, TiC particles tend to show octahedron and/or truncated octahedral shape. The rhomboid and cubic shape of TiC shown in our study is due to the fact that the surface of sample for observation intersects with the TiC crystal with a certain angle.
3.3. Microhardness By taking measurement using Vikers indenter, microhardness was obtained for the laser processed coatings. It demonstrates that high hardness can be imparted to the nano-composite coating (Fig. 8). 3–5 times increase of hardness can be observed for the nano-composite coatings compared with that of the steel substrate. The average hardness values of the coatings are 498 kg/ mm2, 596 kg/mm2, and 689 kg/mm2 for samples processed with 30 mm/s, 60 mm/s, and 90 mm/s respectively. In general, the increase of hardness with increasing laser scanning speed is attributed to the higher content of reinforcements presented in the coating with higher laser processing speed, as indicated by the XRD analysis. Besides the presence of nano-sized TiB2 and TiC
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Fig. 7. (a) Schematic of crystal structure of TiB2; (b) projection of Ti and B atoms along the [0 0 0 1] direction for TiB2 crystal; (c) schematic of crystal structure of TiC; and (d) projection of Ti and C atoms along the [0 0 1] direction for TiC crystal.
improvement of hardness of the nano-composite coating will lead to the increase of wear resistance. Moreover, the in situ formed particles promote strong particle–matrix bonding, thereby reducing pull-out of particles during wear. High toughness is also desirable for materials designed for wear application since it can prevent crack-forming and spallation. It is obvious that the rapid cooling rate associated with laser processing induces the microstructural refinement for both the reinforcements and matrix grains, and reduction of the size of the reinforcements to nanometer size offers the unique opportunity to design wear resistant coating with optimal combination of hardness and toughness [25]. Nevertheless, a comparison of the present study with Ref. [12] shows the overall hardness of TiB2–TiC/Fe is lower than that of TiB2–TiC/Ni. It is attributed to the low volume fraction of reinforcements in the TiB2–TiC/Fe composite coating compared with the TiB2–TiC/Ni, which may result from the complex laser-materials interaction. On the other hand, nickel element plays the role of solid solution to increase the hardness of the matrix in the composite coating. This also gives indication that by introducing alloy elements to the iron matrix through the solid–solution strengthening mechanism the hardness of the composite coating can be further increased.
Fig. 8. Microhardness profiles of the laser processed samples.
reinforcement, the grain refinement effect induced by the a-Fe can also contribute to the increase of hardness. The wear resistance can often be related to hardness through the Archard equation, which states the volume of material worn is inversely proportional to hardness for a given set of test conditions [24]: V ¼K l
P H
ð9Þ
where K is the wear coefficient, l is the sliding distance, P is the applied load and H is the hardness. Thus, it is apparent that the
4. Conclusion TiB2 and TiC reinforced Fe-based nano-composite coating was in situ fabricated by laser surface engineering using Ti, B4C and Fe as the components of precursor. Structural characterization reveals that the composite coating consists of TiB2, TiC, Fe3(B,C) and a-Fe. By comparing integrated intensity of each phase, it is found that the relative amount of TiB2 and TiC increase with the increasing laser scanning speed which was attributed to the
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variation of dilution ratio. Microstructure of the coating indicates that nano-sized TiB2 and TiC particles uniformly distribute in the fine-grained a-Fe matrix. The dual-ceramic phase design is effective in reducing the size of in situ formed reinforcements. Reinforcements of TiB2 and TiC formed following the liquidprecipitation mechanism and their morphology is determined by the crystallographic habit and minimization of surface free energy. Owing to the presence of nano-sized reinforcements (TiB2 and TiC) and the highly refined a-Fe matrix compared with the steel substrate, microhardness value of the composite coating improved significantly (498 kg/mm2, 596 kg/mm2, and 689 kg/mm2 for samples processed with 30 mm/s, 60 mm/s, and 90 mm/s respectively). The preliminary evaluation based on microstructural and hardness analyses show that nano TiB2 and TiC reinforced iron coating using LSE appear to hold promise in wear application. Future work would involve alloying the iron-matrix of the composite coating to further improve its hardness and evaluation of the wear resistance of the nano-composite coating. References [1] Degnan CC, Shipway PH. A comparison of the reciprocating sliding wear behaviour of steel based metal matrix composites processed from selfpropagating high-temperature synthesized Fe–TiC and Fe–TiB2 master alloys. Wear 2002;252:832–41. [2] Zhou Z, Ross IM, Ma L, Rainforth WM, Ehiasarian AP, Hovsepian P. Wear of hydrogen free C/Cr PVD coating against Al2O3 at room temperature. Wear 2011;271:2150–6. [3] Hadad M, Hitzek R, Buergler P, Rohr L, Siegmann S. Wear performance of sandwich structured WC–Co–Cr thermally sprayed coatings using different intermediate layers. Wear 2007;263:691–9. [4] Voevodin AA, Zabinski JS. Nanocomposite and nanostructured tribological materials for space applications. Composites Science and Technology 2005;65: 741–9. [5] Voevodin AA, Hu JJ, Fitz TA, Zabinski JS. Tribological properties of adaptive nanocomposite coatings made of yttria stabilized zirconia and gold. Surface and Coatings Technology 2001;146–147:351–6. [6] Harimkar SP, Dahotre NB. Rapid surface microstructuring of porous alumina ceramic using continuous wave Nd:YAG laser. Journal of Materials Processing Technology 2009;209:4744–9. [7] Harimkar SP, Dahotre NB. Characterization of microstructure in laser surface modified alumina ceramic. Materials Characterization 2008;59:700–7. [8] Agarwal A, Dahotre NB. Comparative wear in titanium diboride coatings on steel using high energy density processes. Wear 2000;240:144–51. [9] Samant AN, Paital SR, Dahotre NB. Process optimization in laser surface structuring of alumina. Journal of Materials Processing Technology 2008;203:498–504.
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