Introduction

Laser cladding is a novel additive manufacturing technique used to fabricate or repair components by depositing clad layers on a substrate. This process uses absorbed laser and is controlled by a CNC system to create a melt pool on the substrate while simultaneously injecting powder into the molten zone to form a deposited track (Ref 1,2,3). Lase cladding is compared to welding process with the ability to produce a strong bond between the substrate and the deposited material. With the provision of good process control, the sharp concentrated and low heat input provides a high precision material deposition (Ref 4). Furthermore, due to its processing advantages such as high-power density and material deposition rates, minimal dilution, low distortion and a strong joint between substrate and deposited material, laser cladding became a popular technique for industrial applications (Ref 5, 6).

High-speed steel (HSS) is considered a high alloy steel and the alloying elements percentage will range between 10 and 25%. These alloying elements will promote carbide formation upon solidification. This steel is presented as Fe-C-X, where X stands for elements leading to the carbide formation such as chromium, molybdenum tungsten and vanadium. HSS is used in applications requiring high strength and hardness, wear and corrosion resistance (Ref 7, 8). During steel solidification, the alloying elements will contribute to the formation of several types of carbides, e.g. MC, M2C, M6C, M7C3 and M23C6 (Ref 9, 10). The carbides type, shape, size, and distribution will directly affect the solidifying HSS mechanical properties. The carbides forming at the grain boundaries will usually have a coarse structure with a size ranging from 20 to 200 µm (Ref 11).

In application where powder injection is used, finer microstructure is expected to appear in the clad layer. Further, the solidification in this process is considered to be rapid, hence leading to superior mechanical properties (Ref 12, 13). It was reported that the M2C will exist in steel containing tungsten and molybdenum, especially with increasing amount of carbon (Ref 14, 15). Though they are metastable at high temperature, during heat treatment this type of carbide will decompose along the γ-austenite matrix into MC and M6C carbides at temperature exceeding 1000 °C (Ref 16). Other studies showed that during the casting of HSS alloys, nano-precipitates evolved and were identified as M7C3 and M23C6 (Ref 17, 18). For ultra-cooled solidifying HSS, as in the case of solidifying liquid melt in laser cladding process, studies showed the presence of an unidentified complex carbide structure rich with V, Mo, W, and Cr. The lattice spacing of the mentioned carbide ranged between 0.15 and 0.26 nm. This spacing is different from the usual reported carbide lattice spacing that tends to exist in HSS (Ref 19,20,21).

As previously presented, many studies addressed the carbide formation for larger specimen, as in casting process, whereas fewer research work addressed carbides in ultra-cooled environment such as in laser cladding where the deposited layer is relatively small and the deposited molten layer will solidify in seconds. Hence this study aimed to provide a comprehensive approach to identify the carbide formation in ultra-cooled T15 HSS. This provides a detailed insight on the microstructure and phase constitution of the cladding layer.

The in-depth characterization of carbide structure and distribution is formulated in different areas of the cladding layer. Initially thermodynamic calculation using equilibrium phase diagram and Scheil–Gulliver model were carried to predict the forming carbides. Then, many experimental procedures to identify the predicted carbides were used such as electron probe microscopy analysis (EPMA), scanning electron microscopy (SEM), and transmission electron microscope (TEM). The characterization process ranged from micro to nano-scale with the addition of identifying the carbides chemical composition.

Materials and Methods

In order to carry the laser cladding experiment, a 3 KW diode laser source (LDM3000-100) was used as it can be seen from Fig. 1. The minimum beam quality was equal to 20 mm mrad and the wavelength range was 900 to 1070 nm. The laser was equipped with a transport fiber having a core diameter of 1000 µm. In order to improve shielding characteristics, the nozzle was equipped by an extra shielding gas tube. Hence, the injection head was equipped with 6 tubes in total, two tubes were for cooling and shielding gas, and the other four were for the powder and the carrier gas.

Fig. 1
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Laserline diode laser source (LDM3000-100)

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The laser beam was directed on the designated substrate in order to create the molten pool. The substrate material was 42CrMo steel. The powder was injected on the molten pool through four circumferentially equally spaced nozzles. The powder material was T15 steel (Febal−x-C-Cox-Cr-Mn-W-V-S), the cladding powder material was purchased with high percentage of carbide forming elements as V and W, the powder had spherical morphology, and the average particle size was equal to 70 µm. The T15 powder material was mixed with CeO2 nano-powder having a mean diameter of 30 nm. The mixture percentage was 99 wt.% of T15 and 1 wt.% of CeO2 powder. The addition of CeO2 into the powder mix aim to deactivate the surface-active elements which are mainly deactivated by Ce. The following will affect the surface tension temperature coefficient resulting outward flow pattern in the molten pool compared to and inward flow for pure T15. This will reduce the segregation at the cladding layer surface.

The carrier gas was argon, the total flux was equal to 6 L/min. Further, argon was also used as a shielding gas, in order to prevent the molten pool from oxidization. The total flux of the shielding gas was equal to 20 L/min. Table 1 presents the material composition and alloying elements of the 42CrMo substrate and T15 powder.

Table 1 Chemical composition of 42CrM0 substrate and T15 powder (wt.%)
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A schematic of the single-track laser cladding experiment is presented in Fig. 2. Fig. 2(a) shows the cladding process with the laser scanning direction, the powder projection onto the substrate and solidifying clad layer. The substrate was static and fixed on the experimental rig. After conducting the experiment, the samples were cut using electrical discharge machining (EDM) where the dimension of the specimen was 1*1 cm including the clad layer. Then they were mechanically grinded and polished with Diamond suspension of 1, 0.5, and 0.25 µm. Finally, the microstructure was examined at transverse and longitudinal cross-sections using TEM and finally XRD analysis as shown in Fig. 2(b).

Fig. 2
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Schematic of laser hot-wire deposition and sample preparation procedure. (a) Laser deposition process (b) Inspected cross sections

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The cut samples were examined under Phenom desktop scanning electron microscope (SEM) in order to characterize the microstructure phases and evolution. Further, electron probe micro-analyzer (EPMA) was used to inspect the elements distribution at different microstructural structure. This experiment was carried using a JFOL jxa EPMA machine.

Microstructure Analysis

Laser trials were carried to produce ideal cladding layer formation and symmetricity. The feed rate was the parameter that controlled the dimension of the clad layer such as height. The laser power and the scanning speed and fee rate were set to be constants and fixed at 600 W, 6 mm/s and 3.6 g/min, respectively.

Microstructure Morphology Characterization

The SEM results of the cut sample cross-sections are shown in Fig. 3. Figure 3 A shows the microstructure at the transverse cross-section and Fig. 3(b) shows it at the longitudinal cross-section. It can be seen that the one clad layer deposited material has good formation quality with absence of cracks and minimal material defects. The interface between the clad layer and the substrate can be clearly seen, further the heat affected zone is also defined as shown below. The substrate–clad interface had circular shape at the transverse cross-section, and linear shape at the longitudinal cross-section.

Fig. 3
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SEM illustration of the cut sample. (a) Transverse cross-section (b) Longitudinal cross section

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The clad layer and substrate magnified microstructure are shown in Fig. 4. Fig. 4(a) and (b) shows a coarse and well-defined as-clad microstructure which can be divided into three phases of microstructure growth. This particular microstructural evolution is due to the high cooling rate of the laser cladding process. At the interface between the clad layer and the substrate, an initial planar growth took place which is due to the presence of the base metal or substrate. The existence of grains at the fusion zone provides the crystal nucleation during solidification, hence new crystals will form while having a certain crystallographic orientation and a cellular growth phase will appear. Finally, an equiaxed stage appeared and dominated the rest of the clad layer microstructure. Due to the ultra-high cooling rate and the relatively small clad layer size, the grains size was almost uniform across the rest of the solidifying layer. Figure 4(c) shows a mgnification of the microstructure at the cellular and equiaxed zones. This revealed the presence of carbide formation at the grain boudary as well as at the austenite matrix. Spherical carbide were present at the inter-granular position (grain boundary) and some on the austenite matrix. Also, needle-like carbides were present at the inter-granular area of the equiaxed zone.

Fig. 4
figure 4

SEM results of microstructure growth patterns and carbides distribution. (a) Transverse cross section (b) Longitudinal cross section (c) cellular and equiaxed zones

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Alloying Material Composition

In order to inspect the alloying material composition in the clad layer, the sample were examined under an electron probe micro-analyzer. The test was repeated several times in order to verify the experiment reliability. The material characterization was carried at 4 points, where the first three points analyzed the material composition of the three stages of growth, at the planar, cellular and equiaxed growth zones. Finally, the fourth point was the inter-granular area at the growth interface of the equiaxed grains, as it can be seen from Fig. 5(a). The inter-granular zone was a subjected of interest because it is expected to have the higher probability for carbide formation, since it will have higher alloying material composition. Furthermore, the EPMA analysis was carried along a line starting from the substrate and passing through the different solidification stages as shown in Fig. 5(b).

Fig. 5
figure 5

EPMA material characterization at different analysis points. (a) (1) planar, (2) cellular, (3) equiaxed, (4) inter-granular (b) material characterization lines

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The alloying materials weight percentage is shown in Table 2, the data were the average of the analysis points repetitive testing and error bars were introduced in Fig. 6 to highlight the measurement differences at each point. Table 2 shows the average weight percentage of all the alloys along with the Fe content. The Fe and Mo mass fraction were not included in Fig. 6 in order to produce a better representation of the alloying material at different inspection points. It can be seen from the graph below that the carbon concentration was around 1.5 wt.% in the planar zone that is roughly around the C wt.% in the HHS powder used in the cladding process. The C wt.% increased to reach 1.76 wt.% in the cellular growth zone, then it increased to reach 2.13 wt.% in the equiaxed zone. Finally, toward total solidification, the C wt.% reached it is highest and was equal to 5.81 wt.% in the inter-granular zone. Along C, other alloying carbides forming elements such as W, Cr and V are presented, as it can be seen from the graph below, W weight percentage was relatively high at points 1,2 and 3, it ranged between 6.34 to 7.43 wt.%, then it reached its highest at point 4 and was equal to 27 wt.%. Similarly, Cr weight increased gradually from point 1 to 3 then it hit 7.89 wt.% at point 4. However, V weight percentage was 2.69 wt.% at point one, then it decreased to 2.41 and 2.57 wt.% at points 2 and 3, then it increased to reach 4.98 wt.% at point 4. The Mn wt.% remained constant through the growth stages and Mo wt.% was low and negligible at points 1, 2 and 3, it reached its highest concentration at point 4.

Table 2 Material weight percentage at the selected points
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Fig. 6
figure 6

EPMA Material weight percentage results

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The alloying elements analysis carried in transverse and longitudinal cross sections aims to present the material distribution trend that will lead to a better understanding on the carbide that might appear at different area of the clad layer. The material distribution showed a carbon weight at points 2 and 3 ranging between 1.76 and 2.13 wt.%, combined with Cr weight ranging between 3.2 and 3.42 wt.% and W weight ranging between 7.11 and 7.43 wt.%. It can be assumed that the carbides which will be present in those areas are identified as MC and M6C rich in the aforementioned alloying elements. As for point 4 that represents the inter-granular zone, there was a high weight concentration of C, which reached 5.8 wt.%, and it had also the highest weight concentration of W equal to 27 wt.% and a weight concentration of Cr equal to 7.89 wt.%, and hence it is assumed that the carbide present in this zone are the M7C3 and M23C6 rich in carbon, chromium and mainly tungsten.

In order to ensure the accuracy and the repeatability of the EPMA test, line material characterization on multiple places in the clad layer was carried and the material distribution trend and uniformity are inspected and the sampling errors were nullified. Hence, it can be seen for Fig. 7 that the material distribution is consistent for each of the characterized material, and the alloying elements C, Cr and W concentration increased as the inspection line traveled from the substrate to the clad layer. On the other hand, the Fe concentration was higher at the substrate and declined in the clad layer.

Fig. 7
figure 7

EPMA signal intensity of C, Fe, Cr and W along the analysis lines presented in Fig. 5(b)

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A closer inspection on the analysis line 1 can be seen from the Fig. 8, initially the content of Fe was high since the Fe wt.% was higher in the substrate than the clad layer. The interface is identified as the doted red line between the substrate and the planar zone. The Fe content was almost constant in the planar zone with some fluctuation peaks in the carbide forming elements Cr and W. As for the cellular zone, the Fe peaks presents the austenite matrix and the C, Cr and W peaks present the inter-granular composition at the cellular grains. Hence the number of Fe peaks can be attributed to the number of grains in the analyzed zone along the line. The Cr and W peaks became more frequent with higher amplitude in the equiaxed zone, and this is due to the grain smaller size compared to cellular zone. The austenite matrix was highlighted by a black enclosing rectangle and the carbides were highlighted by red enclosing rectangles as it can be seen below.

Fig. 8
figure 8

EPMA signal intensity of C, Fe, Cr and W along the analysis line 1 presented in Fig.5(b)

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The Cr and W content peaks recorded in the in the cellular zone indicated the presence of M6C carbides as it can be seen from first highlighted carbide presence in Fig. 8 where the carbon peaked as well; however, its peak was lower in amplitude relative to its total content compared to Cr and W. Cr and W peaks were also present in equiaxed zone with higher intensity. Also, for the second highlighted carbide presence, Cr, W and C peaked in the equiaxed zone with an almost similar intensity which indicated the presence of M7C3 carbide. Finally, the third highlighted carbide presence showed the highest intensity peaks for Cr and W carbides, also C peaked with a relatively lower intensity, this indicated the presence of M23C6 carbide in that area. The MC carbide is hard to detect since it usually precipitates in the austenite matrix and the EPMA analysis line width did not cover enough area on the austenite matrix to detect the presence of the following carbide. It can be concluded that different carbides will be present at different growth stages as it can be seen from the aforementioned data, and as the solidification advanced the alloying elements content will become higher, leading to increased alloying element content in advanced stages of solidification such as in the equiaxed zone. Though the alloying element content increase is slight while the analysis line went from the cellular to equiaxed zone, the increase is noticeable for all alloying elements with the exception of Mn as shown in Fig. 6. Adding that the line analysis is not as accurate as the point analysis, which clearly shows that W content peaked with its highest amplitude at the inter-granular analysis point present in the equiaxed zone. The slight increase between the cellular and equiaxed zone can also be attributed to the short analysis line and the relatively small area where cellular zone is present.

Thermodynamic Calculation of the T15 HSS

Equilibrium Phase Diagram

The HSS Fe-C isopleths were simulated in Thermo-Calc software taking into consideration the alloying elements composition, the diagram is presented in Fig. 9. The carbon composition was equal to 1.6 wt.%, this was highlighted by the red dotted line sketched on the diagram. As it can be seen from the diagram below, the starting precipitation temperature of the austenite phase was 1330 °C and the precipitation temperature of the M6C was 1278 °C. The precipitation temperature of MCI which is presented in Thermal-Calc software as FCC_A1#2 was 1264 °C. The solid stage along with the M6C and MCI carbide started at temperature of 1243 °C.

Fig. 9
figure 9

Isopleth phase diagram of the studied T15 HSS

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The mass fraction variation with respect to temperature is show in Fig. 10. It can be seen that the austenite phase presented by line 5 starting at temperature of 1285 °C. Ferrite phase presented by line 1 started appearing at temperature equal to 798 °C. MCI, M23C6, M6C were present and are presented by lines 2, 3 and 4, respectively. Another M23C6 phase appeared in later stage of cooling presented by line 7. The starting precipitation temperature and maximum mass fraction of carbides started after the decrease in austenite fraction.

Fig. 10
figure 10

HSS phase mass fraction-temperature curve

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The precipitation temperature of MCI carbide was 1274 °C, its mass fraction slightly decreased with temperature decrease. For M23C6 the starting precipitation temperature was equal to 871 °C, the mass fraction increased initially to its maximum then slightly decreased as the temperature decreased. Finally, M6C precipitation temperature was equal to 1292 °C, its mass fraction increased gradually with the temperature decrease. The second M23C6 phase (line 7) precipitation temperature was equal to 383 °C with its mass fraction maximus remaining constant with the temperature decrease.

The main alloying elements site fraction of the three different type carbides with respect to temperature variation are presented in Fig. 11. It can be seen that, for the first precipitated carbide, M6C (Fig. 11a) contained high amount of elements Fe and W and it dissolved some amount of elements Cr and V. As for the second precipitated carbide, the MCI (Fig. 11b) contained high amount of the elements C and V, also it dissolved certain amounts of W, Cr and Fe. Finally, for the last precipitated carbide, the M23C6 (Fig. 11c) contained high amounts of Cr, W and Mo, it also dissolved certain amount Fe.

Fig. 11
figure 11figure 11

Site fraction of carbide forming elements. (a) M6C (b) MCI (c) M23C6

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It can be concluded that the thermodynamic calculation is consistent with the EPMA results carried of the different solidification zone presented in the previous section. At points two and three which presents the cellular and equiaxed growth zones, the EPMA results showed the presence of V, C, Cr and W alloying elements concentration, this led to the assumption of the presence of MCI and M6C carbides, and from the thermodynamic simulation it can be seen that the first precipitated carbide M6C had high amount of W with dissolving amount of V and Cr. Further, for the second precipitated carbide MCI, the thermodynamic simulation showed that it had high amounts of the elements C and V with some dissolving amount of W and Cr which is consistent with the EMPA analysis results. For the EPMA analysis, point 4 representing the inter-granular position at the equiaxed stage, showed high amount of elements W, Cr and Mo. This result was consistent with thermodynamics findings of the M23C6 carbide alloying element content presented in Fig. 11(c), where is showed that this carbide precipitates in later stages of solidification and is rich in the same elements W, Cr, and Mo.

Scheil–Gulliver Simulation

In order to gain a better insight into the carbide formation, and since in laser cladding, high cooling rate is subjected to the clad layer, Scheil–Gulliver model is employed to analyze the carbides evolution in the solidifying layer. Scheil–Gulliver model is often used to study solidifying alloy under non-equilibrium conditions. It provides a more accurate approximation of the solute segregation compared to equilibrium lever rule. Further, it assumes that there is no diffusion in the solid phase in order for the liquid composition to follow the liquidus temperature represented by the equilibrium phase diagram. Figure 12(a) presents the molar fraction of solid phases and the precipitation temperature of each phase. Also, Fig. 12(b) shows the weight fraction of each phase. It can be seen from Fig. 12(a) the presence of the following carbides: M6C, MCI, M7C3 and MCII referred as MC_SHP in thermal-Calc.

Fig. 12
figure 12

Scheil–Gulliver model Mole and weight fractions variation with respect to temperature. (a) Mole Fraction (b) Weight fraction

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As it can be seen for M6C and MCI the precipitation temperatures were 1278 and 1264 °C respectively, this was similar to the precipitation temperature of those carbide calculated in the equilibrium phase diagram from the previous section. Two new carbides appeared in Scheil–Gulliver simulation model, the M7C3 and MCII carbides having the same precipitation temperature equal to 1192 °C.

When compared to the equilibrium phase diagram, Scheil–Gulliver model also predicted the precipitation of MCI and M6C type carbides that will appear at early stages of solidification. Furthermore, two new carbides appeared which were M7C3 and MCII type carbides, the precipitated shortly after the appearance of MCI and M6C. This is due to the Scheil–Gulliver model properties assuming that the solid stage has no diffusion, hence as the solidification proceeds, the alloying elements will not precipitate in the solid part. This will increase the content of the alloying elements in the liquid phase, hence M7C3 and MCII carbides rich in elements such as W and Cr appeared. As shown from the equilibrium phase diagram before, M23C6 will precipitate in later stages of solidification, hence this carbide was absent from the Scheil–Gulliver simulation due to the fact that after total solidification, the diffusion will not proceed in the solid stage.

Discussion

Figure 13 is a representative schematic of microstructure resulting from the laser cladding of HSS studied in this paper. The presented microstructures were the planar, cellular and the equiaxed zones. As it can be seen, the carbides formed at the inter-granular zones and some were present at the austenite matrix. Granular and needle like carbides at the grain boundaries and in the austenite matrix were identified. Table 3 shows the precipitation temperature of the carbides identified by the equilibrium phase diagram and the Scheil–Gulliver model. The starting microstructure consisted of small granular M6C carbide at the grain boundaries of the cellular zone as it can be seen from Fig. 13 the following can be concluded from the EPMA and site fraction results showing that area is rich in W and Cr and from the fact that M6C had the highest precipitation temperature.

Fig. 13
figure 13

Illustration of the different solidification zones with their respective carbides structure and distribution

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Table 3 Carbides precipitation temperature
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The second precipitated carbide was the MCI, having a slightly lower precipitation temperature compared to M6C, the following carbide is known to precipitate in the austenite matrix (Ref 22). Further, it is rich in C and V which was shown in the analysis of points 2 and 3 of the EPMA and from the equilibrium phase diagram site fraction results. The MCI carbide is highlighted in figure below.

The last precipitating carbides were identified as M7C3 and M23C6 which appeared at the grain boundaries of the equiaxed zone. They hold the lowest precipitation temperature as shown in Table 3. The presence of M23C6 was predicted form the equilibrium phase diagram analysis and Scheil–Gulliver model predicted the presence of M7C3. The M7C3 and M23C6 carbides are rich in Cr, W and C. Hence, the EPMA results at point 4 showed the presence of those carbides. As it can be seen in Fig. 13 the granular smaller carbides were identified as M23C6 as they tend to form as small spherical like grains along the grain boundary (Ref 23, 24). Also, usually carbides will form a cluster of different type of carbides, hence clusters of carbides can be seen for M6C and M23C6. Finally, the M7C3 was identified as needle-like carbides which were dispersed along the grain boundary and also formed clusters at some certain areas (Ref 24, 25).

Scheil–Gulliver model predicted the precipitation of MCII, however, as it can be noticed from the weight fraction of each carbide presented in Fig. 12, its weight fraction in the solid phase is negligible. Adding that it is hard to index smaller amount of carbides in the EPMA and SEM analysis, hence the following carbide was not identified.

In order to verify the predicted carbides that will be present in the clad layer after the laser cladding process, scanning transmission electron microscopy (TEM) is carried and the carbides morphology and selected area diffraction patterns (SADP) are presented in Fig. 14. The carbide presented in Fig. 14a is identified as MCI, which will be present on the inside of the Fe matrix. The carbide presented in Fig. 14(b) is identified as M23C6 having a relatively small size and round shape spawning at the grain boundary. The carbide cluster shown in Fig. 14(c) is identified as M7C3 and finally the carbide presented in Fig. 14(d) is identified as M6C which appeared at the grain boundary. As anticipated by the EPMA material characterization and phase diagram analysis, the TEM verified the presence of the predicted carbides in the Fe matrix and at the grain boundary.

Fig. 14
figure 14

Fine precipitates after laser cladding: (a) MCI in a bright field (BF) image from scanning transmission electron microscopy (TEM), (b) M23C6 in a TEM BF image, (c) M7C3 in a TEM dark field (DF) image, (d) M6C in a TEM DF, (e) SADP of MCI in (a), (f) SADP of M23C6 in (b), (g) SADP of M7C3 in (c), and (h) SADP of M6C in (d)

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Figure 15 shows the XRD results of the deposited track. The main precipitated phases were MCI, M6C and M23C6 carbides. The XRD results did not detect the presence of M7C3 and MCII carbides, this is can be attributed to their low weigh fraction as it can be seen from the Scheil–Gulliver model simulation presented in Fig. 12.

Fig. 15
figure 15

Laser cladding deposited track XRD results

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Conclusion

The microstructure of T15 high-speed steel used in the laser cladding experiment was studied and characterized. The different alloying elements at different microstructural phases were inspected and the carbides types were predicted. The relationship between the microstructural phase and carbides alloying elements content was established. The alloying elements evolution starting from the substrate and passing over the clad layer was recorded and analyzed. The thermodynamic simulation on the T15 HSS was accomplished and the predicted carbide types were verified. Several conclusions can be drawn as follows:

  1. 1.

    The alloying element analysis in the cellular and equiaxed zones showed relatively high concentration of C, Cr, W and V which indicated the presence of MCI and M6C carbides. Similarly, the alloying element analysis at the inter-granular zone showed the highest concentration in W and Cr which indicated the presence M7C3 of M23C6 carbides.

  2. 2.

    The alloying elements characterization along the line starting from the substrate and passing through the clad layer showed the presence of carbides alloying elements in the austenite matrix and the inter-granular zones. Also, it indicated an increase in alloying element content as the analysis line hit the interface and proceeded into the clad layer, where the amplitude peaks hit their highest in the inter-granular zone.

  3. 3.

    HSS equilibrium phase diagram showed three type of carbides which precipitated respectively and were M6C, MCI and M23C6. The elements site fraction of the carbides showed that M6C contained a relatively high amount of W and Fe, the MCI carbide had high content of C and V, and the M23C6 carbide contained the highest content in Cr, W, and Mo. The thermodynamic simulation verified the carbides predicted in the previous experimental section, with the only carbide absent from the equilibrium phase diagram being the M7C3.

  4. 4.

    The Scheil–Gulliver simulation model revealed the presence of four carbides type identified as M6C, MCI, M7C3 and MCII which precipitated, respectively, with the decrease in temperature. Similar to the equilibrium phase diagram, Scheil–Gulliver model predicted the precipitation of M6C and MCI. However, it predicted the precipitation of the two additional carbides M7C3 and MCII. Further, this model indicated that the M23C6 will not precipitate, this is due to the assumption of the Scheil–Gulliver that there is no diffusion in the solid state, and the M23C6 will precipitate the latest as shown in the equilibrium phase diagram.

  5. 5.

    TEM results showed that the predicted carbides were present in the clad layer where all the carbides are identified with the exception of MCII carbide. Similarly, XRD results predicted the presence of MCI, M6C and M23C6 carbides, however M7C3 and MCII were not detected due to their low weigh fraction.