Laser cladding of SiC multilayers for diamond deposition on steel substrates

Diamond & Related Materials 65 (2016) 105–114

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Laser cladding of SiC multilayers for diamond deposition on steel substrates Andre Contin a,⁎, Getúlio de Vasconcelos b, Danilo Maciel Barquete c, Raonei Alves Campos d, Vladimir Jesus Trava-Airoldi a, Evaldo José Corat a a

National Institute for Space Research — INPE, Av. dos Astronautas, 1758, Jd. Granja, 12227-010 São José dos Campos, SP, Brazil Institute for Advanced Studies, — IEAv, Trevo Coronel Aviador José Alberto Albano do Amarante, 01, Putim, 12228-001 São José dos Campos, SP, Brazil State University of Santa Cruz — UESC, Rod. Jorge Amado, km 16, Salobrinho, 45662-900 Ilhéus, BA, Brazil d Federal Institute of Education, Science and Technology of São Paulo Brazil – IFSP, R. Antônio Fogaça de Almeida, S/N, Jardim Elza Maria, 12322-030 Jacareí, SP, Brazil b c

a r t i c l e

i n f o

Article history: Received 30 November 2015 Received in revised form 5 February 2016 Accepted 5 February 2016 Available online xxxx Keywords: Diamond film Laser cladding HFCVD

a b s t r a c t It is well-known that growth and adhesion of polycrystalline diamond coating directly on steel are both problematic. To solve these issues, interlayers are needed. For the present study, diamond film was deposited on steel with intermediate barrier of silicon carbide (SiC). In addition, laser cladding process produced interlayer. Diamond films were grown by Hot Filament Chemical Vapor Deposition (HFCVD). During laser cladding process, SiC was partially dissociated and so, formation of FeSi was observed. Just 1-layer SiC was enough to grow diamond film since the iron is no longer free to migrate to the surface during the CVD deposition. Nevertheless, the high residual stress formed at the interface, during the reactor cooling, produces fragmentation of the diamond film grown with 1-layer SiC on the steel. To overcome the disadvantage, we propose to create SiC multilayers in an effort to reduce the influence of steel thermal expansion coefficient. Detailed characterization of the SiC interlayer, FeSi phase and diamond coating, are discussed based on X-ray Diffraction, Scanning Electron Microscopy and Raman Spectroscopy. Results showed that the presence of FeSi, formed by SiC dissociation, was the main reason for diffusion barrier effectiveness. Further, the use of SiC multilayers efficiently relaxed the high thermal stress. © 2016 Elsevier B.V. All rights reserved.

1. Introduction In general, alloy steels are the most employed in modern industry; however, there is a drop of its useful life when applied for critical wear resistance. Hard carbon coating, such as diamond is a possible solution to minimize this issue. CVD diamond coating is an excellent candidate for several applications, because of its properties, such as high mechanical hardness, high thermal conductivity, high electrical resistivity, low friction coefficient and corrosion resistance [1]. Thus, diamond coating on steel may improve its hardness and wear resistance. It is a desirable coating for industrial applications, mostly in machining applications. Despite huge industrial potential, direct diamond coating on steel is still not effective because of poor adherence between diamond and steel caused by three major issues. Iron has a high vapor pressure and catalyzes graphite soot formation at interface. Because of high carbon solubility in iron at CVD temperature, carbon diffusion inward steel decreases the diamond nucleation rate. The high difference in thermal

⁎ Corresponding author. E-mail address: [email protected] (A. Contin).

http://dx.doi.org/10.1016/j.diamond.2016.02.007 0925-9635/© 2016 Elsevier B.V. All rights reserved.

expansion coefficients between diamond and steel, a mismatch that induces high stress within coated diamond [2,3]. One way to solve these problems is an intermediate layer, which acts as diffusion barriers, for both iron and carbon. A series of works has revealed the use of diffusion barriers including iron boriding [3], chromium nitride [4,5], aluminum coatings [6], silicon film [7], titanium nitride [8], tungsten [9] and titanium carbide [10]. Most of these works efficiently blocks iron and carbon diffusion to enable good quality diamond growth but do not get rid of the high stress because of thermal expansion mismatch. The Fe2B alternative by Buijnsters et al. [3] got low stress, but film quality was limited. Polini et al. [10] showed a relation between interlayer roughness and diamond coating thickness on both stress development and adhesion. The laser cladding produces surface modification. It creates a dense coating, with laser irradiation on a powder dispersed on substrate surface, and a strong metallurgical bond with substrate [11]. The advantages of laser cladding over conventional techniques to create a diffusion barrier include a faster processing speed, high heating/cooling rate, precision, automation and versatility. To our best knowledge, there is no literature in using laser cladding to form interfaces for CVD diamond deposition. This paper reports on diamond growth on a 304 stainless steel coated with laser cladding for creating a SiC intermediate layer. Laser

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cladding employs a high power laser beam to melt quickly the powder material and then rapidly solidify to create a coating that is bonded to the substrate [12]. The SiC has more than 200 polytipes and is distinguished by the stacking sequence of each tetrahedrally bonded Si\\C bilayer [13]. The SiC is inert to the gas atmosphere used for diamond deposition and has an intermediate thermal expansion coefficient; it is a good candidate interlayer material for diamond coating on steel [14,15]. This study explores three thicknesses of intermediate layers and their influence on diamond coating. 2. Experimental details The substrate for diamond deposition was a commercial 304 stainless steel cut into disks of 25.4 mm diameter with 3 mm thickness and mechanically polished with 320 and 600# sandpapers. The radiation source for laser cladding was a carbon dioxide laser (Synrad-SH) with output power of 125 W and beam diameter of 200 μm. Beam intensity was 400 kW/cm2. Nitrogen flow provided environment purging to avoid oxidation. The powder to create the intermediate barrier was a commercially available silicon carbide powder (Treibacher Schleifmittel Brasil) with predominance of rhombohedral crystal system. SiC average grain size was 4 μm in all experiments. We used two-step process (pre-placed laser cladding). In this process, the first stage consists of powder application by air gun. The second stage is laser treatment. For 3- and 6-layer SiC, both the powder pulverization and laser irradiation were repeated 3 and 6 times, respectively. The laser parameters are described in Table 1. These laser parameters were chosen because they created favorable results in preliminary experiments. Before diamond deposition, we performed substrate seeding. The diamond nanoparticle seeding was an Electrostatic Self-Assembly (ESA) process [16] consisting of two steps. The first step was sample immersion for 30 min in an aqueous PDDA solution {-Poly (diallyldimethylamonium chloride)}. The second was immersion in aqueous solution containing 4 nm dispersed diamond powders and the anionic polymer PSS-(Poly (sodium 4-styrenesulfonate), for a 30 min period. After each step, we washed the substrates in DI water [17]. Diamond film deposition took place in a HFCVD (Hot Filament Chemical Vapor Deposition) reactor. A set of five 125 μm tungsten filaments, kept at a temperature around 2200 °C and positioned at 5 mm above the substrates was the activation region. The total gas pressure was 50 Torr during the 3 h growth. The reactive gas mixture consisted of 2% CH4 and 98% H2. The ramp down turn off period was of 1 h. Under these CVD conditions, the deposition rate, at 700 °C, was approximately 0.43 μm/h. The deposition rate at 550 °C was approximately 0.066 μm/h. Characterization of samples included Field Emission Gun-Scanning Electron Microscopy (FEG-SEM Mira 3 Tescan), X-ray diffraction (XRD Panalytical X'pert Pro using CuKα radiation at wavelength of 1.5418 Å) and Raman Scattering Spectroscopy (Horiba Scientific) with a helium cadmium laser excitation (325 nm). We examined diffusion bonding at the interface by EDX linescan (Oxford max 50). The Reicherter (BVR-1083) Rockwell Hardness Tester with diamond 120° cone angle evaluated laser cladding layer adhesion. Surface roughness was determined by FEG-SEM images with the MEX software from Alicona. Table 1 Laser cladding parameters. Laser parameters

1-layer SiC

3-layer SiC

6-layer SiC

Resolution (dots per square inch) Number of heating cycles Beam scanning speed (mm/s) Atmosphere

600

600

600

1 300

1 300

1 100

N2 − 5 l/min flow

N2 − 5 l/min flow

N2 − 5 l/min flow

3. Results and discussion Fig. 1(a,b), (c,d) and (e,f) shows the SEM top view image of 1-,3- and 6- layer SiC, respectively, deposited by laser cladding on the 304 stainless steel substrates. A close comparison, at higher magnification, among Fig. 1(b, d and f) shows significant differences in morphology. Fig. 1(a,b) shows the remelting process, with high diffusion of elements, SiC particles can't be observed on matrix. Fig. 1(c,d) shows SiC sintered on the surface, forming sintered crusts in some areas. The 6-layer SiC sample (Fig. 1(e,f)) shows granular structure on surface. There was only some SiC powder coalescence and neck formation featuring low powder densification. The skeleton-like structure points out weak bond between the particles. Fig. 2(a,), (b) and (c) shows the X-ray diffractogram of one 1-, 3- and 6- layer SiC, respectively, on the substrates. Therefore, the 1-layer SiC samples have austenite and FeSi highest peak intensity, showing SiC dissociation during laser irradiation. For 3-layer SiC samples the FeSi content increases, confirming the SiC dissociation, and there is some unreacted SiC. With 6-layer SiC austenite peaks almost disappeared because of the increase of SiC thickness, but there is still FeSi alloy among SiC grains. Majumdar et al. also found a SiC partial dissociation during the laser irradiation. In their work, there wasn't also the formation of cementite (Fe3C) by the extra carbon reaction of SiC, which is associated with fast heating and cooling of laser; the formation of the cementite was suppressed at low temperatures during the ramp down [18]. One of the possible mechanisms of FeSi phase formation is silicon collision with iron in the molten pool during laser irradiation [19]. The FeSi phase is intermediate between the iron-rich phase (Fe3Si) and silicon-rich phase (FeSi2). The FeSi phase appears with diffusion around 50 at.% of silicon, [20]. According to Buijnsters et al. FeSi phase is efficient as diffusion barrier for Fe. Apart from that Fe diffusivity in FeSi is in order of b10−14 m2/s, which is lower than Fe diffusivity in Fe3Si phase [7]. Fig. 3 shows the sample interface cross section as observed by FEG-SEM images. For a 1-layer SiC (Fig. 3(a)) it is not distinguishable the location of a defined interface between the steel and the SiC layer. Fig. 3(b) shows, for 3-layer SiC, the formation of intermediate layer with sintered granular structure. Fig. 3(c) shows similar result for 6-layer SiC. Thickness measurement was performed by the EDS linescan shown in Fig. 4. Fig. 4 shows EDS linescan from cross sections. For 1-layer SiC (Fig. 4(a)) there was a total diffusion of element inward steel. Both, 3- and 6-layer SiC show a dilute interface between steel and SiC because of interdiffusion process. It is important to note the high silicon diffusion into steel, above 25 μm, as shown in Fig. 4(b and c). High silicon thermal diffusion, at the first SiC layer on steel, ensures an excellent metallurgical bond without any visible porosity. A strong bond between the coating and substrate is required because of huge compressive stress induced on SiC layer during the reactor cooling, mostly by steel high thermal expansion coefficient. SiC coating thickness was around 12 μm for 3-layer SiC and 18 μm for 6-layer SiC. The thickness is not constant along the cross section of the samples since clad has high roughness because of particles points out weak bonded to the surface and overlapping tracks. Thickness does not double from 3- to 6-layer SiC because of increasing roughness. Probably part of the sprayed SiC particles of one layer cover up voids left in previous layer. About 8 and 6 μm depth from the substrate surface, as shown in Fig. 4(b) and (c) respectively, it is possible to view a small amount of iron migrated from bulk material. This iron migration among SiC particles was responsible by extension of FeSi phase formation, which probably acted as a binder among these particles. Fig. 5 shows the surface micrographs at region indented with 600 N load on stainless steel with 3- and 6-layer SiC.

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Fig. 1. SEM top view image of SiC deposition by laser cladding of: (a,b) 1-layer SiC, (c,d) 3-layer SiC and (e,f) 6-layer SiC.

Fig. 5(a) displays no delamination and cracking around the imprint. Fig. 5(b) shows the indent area with spalling and partial flaking of SiC layer. This figure clearly shows the presence of radial cracks running from the indent. In brittle layers, the contact stress field creates this kind of indentation fracture [21]. Obviously the six SiC layer samples are more fragile.

Although the laser parameters are not enough to create of total melt pool in the top surface of pre-placed powder, they can sinter the SiC particles. Heat absorbed by the powder layer transfers to the substrate by conduction. The heat transfer, if efficient, results in a molten pool formation on substrate surface. This melt wets SiC particles, probably creating conditions for Fe and Si alloying. Probably, for the 1-layer SiC

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Fig. 2. X-Ray diffraction pattern of (a) 1-layer SiC, (b) 3-layer SiC and (c) 6-layer SiC.

the full alloying occurred because of the powder full immersion in molten substrate. Up to three layers there was enough heat transfer to mix up some molten iron among SiC layers and to form a FeSi binder. This results in a less fragile interlayer (see Figs. 1(c,d) and 5(a)). In contrast, when number of layers increased, the heat transfer to the substrate was not enough to melt substrate surface and form a binder.

Fig. 3. Cross section of (a) 1-layer SiC, (b) 3-layer SiC and (c) 6-layer SiC.

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Fig. 4. Linescan of the cross section of (a) 1-layer SiC, (b) 3-layer SiC and (c) 6-layer SiC.

This condition limited SiC particles sintering which formed the brittle layer. In conclusion, small amount of free iron is essential as a bind material to reduce porosity and increase layer densification. From the results obtained, the fused layer provides the requirements to act as an intermediate barrier for diamond deposition. Fig. 6(a,b) shows diamond film grown on steel with 1-layer SiC. As can be seen, diamond coating has fragmented into small pieces. The explanation for diamond film fragmentation is the high thermal stress induced by the substrate during reactor cooling down. The

diamond failure mechanism on substrate coated with SiC layer is different from bare steel. The fragmentation into tiny pieces indicates a film tightly adhered to substrate and it's collapsing because of the high thermal stress during cooling down. On bare steel diamond grows over a loose graphitic interlayer and disrupts under free-standing state with large diamond pieces delamination during ramp down cooling. [22]. 1-layer SiC was unable to minimize the thermal stress at diamond interface, but it was enough to grow good quality

Fig. 5. Rockwell indentation test of (a) 3-layer SiC and (b) 6-layer SiC.

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Fig. 6. FEG-SEM image of diamond film on steel of (a,b) 1-layer SiC, (c,d) 3-layer SiC and (e,f) 6-layer SiC.

diamond film. The FeSi alloy developed acted as diffusion barrier so that iron was no longer free to migrate to the surface during the CVD deposition. Fig. 6(c,d) and (e,f) show the diamond films deposited on samples with 3- and 6-layer SiC, respectively. In contrast to the diamond film obtained at 1-layer SiC the diamond coatings on 3 and 6 layers are completely continuous and uniform. With 3- and 6- layer SiC, almost

all cast crusts overlap previous clad, some connects to previous, and others attaches within intervening areas. It is obvious that the laser cladding forms intermediate layers with high surface roughness, which increases the diamond film adhesion by mechanical interlocking [23]. Laser irradiation creates high surface roughness, and during early deposition stage, the diamond nuclei grow inside the surface cavities of SiC layer, producing the interlocking action.

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Fig. 7. X-Ray diffraction pattern of 3-layer SiC after CVD diamond growth.

The surface topography at higher magnification in Fig. 6(b), (d) and (f) shows well faceted polycrystalline diamond films, no porous graphitic are observed on the top surface. The grainy structure of the microcrystalline diamond films shows individual and groups of faceted diamond grains with both {111} and {100} surface morphology. The Fig. 7 shows the X-ray diffractogram of 3-layer SiC after CVD diamond growth. In this sample, diamond film fragmented on the surface and it was removed from the substrate. From the diffractogram, it can be inferred that FeSi phase kept higher peak intensity after diamond growth. The appearance of single Fe1.34Si0.66 peak might be attributed to the minimum iron diffusion inward FeSi which caused an increase in Fe content, in the phase. On the other hand, the diffractogram shows the silicon-rich phase (FeSi2) appearance. Even with surface temperature at 700 °C, no graphite peak is seen, despite the tendency for iron diffusion. Although, iron silicides are not efficient as diffusion barrier for carbon [7], the diffractogram shows a unique peak related iron carbide formed at the surface during the CVD growth. Carbon diffusivity does not decline the diamond film quality, but it can delay the diamond nucleation [22]. On the other hand, we used the diamond nucleation process based on ESA. This method creates a monolayer of diamond nanoparticles on functionalized substrate. Thus, seeding process embeds abundant diamond particles on surface substrate, improving the diamond nucleation rate [17]. Fig. 8(a) shows the Raman spectrum of diamond film on 1-layer SiC excited at 325 nm. The sharp peak centered at 1334 cm−1 corresponds to the diamond lattice vibration. The band observed at 1590 cm−1 is correlated to sp2 amorphous carbon, known as G-band, while the band at 1380 cm−1, so called D-band of graphite, corresponds to sp2 disorder [24]. Although the film features regions where delamination took place, other adjacent zones are adhered. In this area, even with residual stress relaxation around it, the spectrum exhibits the displacement of 2 cm−1 from the diamond characteristic peak. This is correlated with good adhesion at interface. This fact reinforces that the film delamination was induced by high thermal stress at interface and not by lack of adherence. Fig. 8(b) and (c) show the relatively sharp peak centered at 1334 cm−1 and 1333 cm−1 respectively, characteristics of diamond structure. At 1590 cm−1 is detected carbon sp2 hybridized. The spectra showed high purity of diamond because of low intensity of G-band, matching to a low graphitic phase. FeSi phase proved to be an efficient barrier for iron diffusion from the steel. Diamond film grown on rough surface tends to reproduce the surface topology until the film thickness exceeds the surface roughness

Fig. 8. Raman spectrum of diamond film on steel of (a) 1-layer SiC, (b) 3-layer SiC and (c) 6-layer SiC.

value. In this case, CVD growth parameters define the surface topology, instead of underling layer roughness [25]. The surface roughness values of diamond film (see Fig. 6(c)) covered on the peak region of 3-layer SiC were Ra = 2.82 μm, Rq = 3.15 μm and Rz = 3.96 μm. For the flat region, the values were Ra = 1.85 μm, Rq = 2.16 μm and Rz = 2.91 μm. For diamond film (see Fig. 6(e)) covered the peak region of 6-layer SiC, the values were Ra = 2.07 μm, Rq = 2.48 μm and Rz = 3.30 μm.

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Fig. 9. FEG-SEM image of diamond film on steel of (a,b) 1-layer SiC at 550 °C.

For the flat region, the values were Ra = 1.81 μm, Rq = 2.30 μm and Rz = 3.38 μm. Roughness analysis on diamond films of 3- and 6-layer SiC showed the surface roughness value (peaks and flats regions) greater than diamond film thickness (1.3 μm). Under these conditions, diamond spatial position on rough layer and film thickness influence the stress values of diamond film [10]. The spectra of Fig. 8(b) and (c) were collected at peak regions covered by diamond coating. The diamond peaks were centered at 1336 and 1337 cm−1 on flat regions for 3- and 6-layer SiC, respectively. It is possible to suggest that local of peak region reduced shrinkage of stress induced by steel substrate. Also, the silicon carbide has intermediate thermal expansion coefficient between diamond and steel, this fact accommodated the high steel thermal stress along of the intermediate barrier. In spite of stress relief contribution, a SiC interlayer constitutes a brittle intermediate coating that can compromise the mechanical properties of the diamond film. In order to decrease the thermal stress at interface, diamond film was grown on 1-layer SiC with surface temperature at 550 °C, the other deposition parameters were the same. Fig. 9 shows the film deposited at 550 °C.

Fig. 10. Raman spectrum of diamond film on steel of 1-layer SiC at 550 °C.

The deposited film is continuous and adherent without fragmentation, the image at higher magnification (Fig. 9(b)) reveals clusters of diamond grains in nanometric scale, diamond crystals present ballaslike structure with spherical morphology. The film grown on 1-layer SiC at 550 °C has smaller thickness than films at 700 °C. Smaller thickness induces a lower stress. In addition, low temperature reduces grain size, which also lessens the stress. Therefore, minimal stress value is obtained for low deposition temperature [26]. Fig. 10 shows the Raman spectrum of diamond film grown with surface temperature at 550 °C. The spectrum exhibits shoulders at 1150 and 1490 cm−1 together, indicating the presence of transpolyacetylene (TPA) at the grain boundaries [27]. The band at 1338 cm−1 is a characteristic of carbon sp3 hybridized, shifted of about 6 cm− 1 from the diamond characteristic peak. The magnitude of biaxial stress into the film is relative to degree of diamond peak shift [28]. Therefore, Raman features in Fig. 10 the formation of nanocrystalline diamond [16]. Though the iron diffusion effect on the diamond coating quality is likely to decrease with decreasing temperature, there will be a decrease H-atom abstraction rate, by leading to greater amount of amorphous carbon [7]. The surface roughness values of diamond film (see Fig. 9(a)) of 1-layer SiC were Ra = 347.72 nm, Rq = 436.66 nm and Rz = 1.03 μm. Because of dissolution/diffusion of Si and C into steel and deposition of grains in nanometric scale, the surface roughness value of diamond on 1-layer SiC is lower than diamond on 3- and 6-layer SiC. On the other hand, the roughness is still greater than film thickness (200 nm) since growth rate reduced at lower temperature. Therefore, surface roughness greater than film thickness and low deposition temperature contributed to the deposition of adherent diamond films on steel using 1-layer SiC. It is recognized that SiC nanowire structures can be synthesized through heating carbon and silicon precursors in the presence of iron [29,30]. The substrate containing Fe (i.e., steel) can be a source for Fe vapor. Zhuang et al. [31] developed a method to create SiC nanowire, which uses an iron source as catalyst and nanodiamonds (5 nm) as diamond seeding, to produce SiC nanocable/diamond hybrid structure on SiO2 substrates by MWCVD process. In this work, the iron is no longer free to migrate to the surface during diamond growth to act as catalyst for SiC nanowire formation. The FEG-SEM images of diamond film (Figs. 6 and 9) do not show the SiC nanowire structures grown on diamond film [31]. In addition, from Raman analysis (Figs. 8 and 10), it is not observed two peaks (780 and 951 cm − 1 ) related to SiC nanowire structures [30].

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Altogether, these results showed that silicon carbide interlayer deposited by laser cladding on steel substrates changed the surface conditions for diamond deposition. Inward silicon diffusion and creation of FeSi alloy suppressed graphite soot formation on the steel. This provided carbon source for diamond nucleation. The thick FeSi layer created by remelting process, with just 1-layer SiC, was enough to grow diamond film since the iron was no longer free to migrate to the surface during the CVD deposition. Raman analysis showed good quality of diamond film. Nevertheless, the high residual stress formed at the interface, during the reactor cooling, produced fragmentation of the diamond film grown with 1-layer SiC on steel. The growth on samples with 1-layer SiC at lower temperature (550 °C) reduced the residual stress and showed total adherence of diamond film on the steel. 3- and 6-layer SiC were tried to decrease thermal stress even at 700 °C diamond deposition temperature. The morphology of 3-layer SiC produced SiC sintered particles on the surface. In contrast, 6-layer SiC showed weakly sintered particles, which resulted in a brittle interface. In all cases, Fe and Si alloying promoted a good metallurgical bonding of SiC interlayer with substrate. Diamond films deposited on 3- and 6-layer SiC at 700 °C, in contrast to film on 1-layer SiC, remained adhered and continuous after cooling down. Besides the relaxation of residual stress along SiC layers, the high rough layer allowed diamond film stability. A brittle interlayer can compromise diamond film anchoring if demanded mechanically. A possible solution would be the use of a binder material among SiC powder. A powder of this binder could be mixed with SiC one, but a preferable approach is to deposit alternate binder and SiC layers. Which binder material is of choice is under current study in our laboratory. The Fe and Si alloying showed above indicated that even iron could be a binder. However, other materials are under test. A binder could also decrease surface roughness left by the only partial sintering of SiC particles. Further work is in progress to create a composite layer, using SiC particles and binder material, like titanium or cooper powder.

4. Conclusions In conclusion, we have shown that SiC layers deposited by laser cladding fulfilled the conditions for diamond deposition on ferrous substrates. The coating showed good metallurgical bonding with steel and leaded to a good diffusion barrier. This barrier avoided graphitic formation. The layers proved to be also effective in limiting carbon diffusion during the diamond deposition. Results suggested FeSi, formed by SiC dissociation, as the main reason for diffusion barrier effectiveness. The 1-layer SiC experiments revealed that Si alloying modified steel surface and allowed direct diamond growth. Further, the use of 3- and 6-layer SiC efficiently relaxed the high thermal stress produced by the huge difference between steel and diamond thermal expansion coefficients. Further, this study introduces laser cladding as a worthy method to create interlayer for CVD diamond growth.

Prime novelty statement Laser cladding proves to be the worthy method to create intermediate barrier for diamond deposition on ferrous materials, the feasibility of using SiC interlayer produces diamond films adherents to the substrate.

Acknowledgments The authors would like to thank FAPESP (Process Number 2013/25939-8) for the financial support of this work.

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