Laser Surface Remelting in Single-Crystal Nickel-based Superalloy using a Continuous Wave Fiber Laser

Single-crystal (SC) nickel-based superalloy castings offer high-temperature microstructural stability and superior creep resistance, due to which they are exclusively used in the hot sections of gas turbine engines. However, SC nickel-base superalloy components are ‘difficult-to-cut’ while manufacturing. Worldwide research shows an interest in improving the machinability of superalloys. The present work investigates the controlled surface damage on CMSX-4 superalloy through laser surface remelting technique towards the improvement of their machinability. The specimens were laser-treated using a constant laser power and scan speed and by varying the positive focal position to get a range of energy densities. The process, structure, and property were systematically studied in the fusion zone (FZ). The FZ shape changes from keyhole to conduction mode with the increasing focal position. The FZ shows a finer assorted dendritic structure and lesser elemental segregation than the base metal (BM). In keyhole mode penetration, large pores, and multi-directional cracks were observed in the base region. On the other hand, the conductive mode showed only vertical centreline cracks and no significant porosity. The cracks are attributed to thermal stresses and elemental segregation produced during solidification. Microcracking was also observed near the fusion boundary and is attributed to the presence of low melting Mo and Ti–rich eutectics. The FZ away from the cracks showed 10% lower hardness than the BM, which is attributed to the dissolution of γ′ phase. Overall, the laser processing under the given range of energy densities produced wide variants of surface defects in the FZ.


Introduction
Continuous wave (CW) solid-state fiber lasers are extensively used in the processing of advanced materials due to their high beam quality, brightness high energy efficiency, and portability [1,2]. Nickel-based superalloys, owing to their stable high-temperature strength, high creep resistance, and resistance to corrosion and oxidation, are exclusively used in the hot sections of advanced gas turbine engines [3,4]. When cast as single-crystals (SC), the absence of transverse grain boundaries enhances their creep resistance [5]. Due to their high strength, significant work hardening, and low thermal conductivity, these SC nickel-based superalloys become difficult to machine while manufacturing turbine blades [6]. The way to reduce cutting forces is to introduce defects in the surface layers. Efforts to introduce controlled damage in the surface layers of SC nickel-based superalloy CMSX-4 using CW fiber laser have been reported by the present authors in a recent publication [7]. The use of different energy densities in laser surface melting can result in two different modes of penetration: conductive mode and keyhole mode. The energy density in fiber lasers can be regulated by adjusting the laser power, scan speed, and beam diameter. In the reported work, a constant laser beam diameter of 1.2 mm and laser power of 1000 W with varying scan speeds from 5.5 mm/s to 16.7 mm/s was used while maintaining the energy density in the range of 50 to 150 J/mm 2 , which resulted in 'conductive' mode of heating, producing shallow and wide fusion zone (FZ). The FZ showed the presence of various defects such as microstructural changes, solidification cracks, and microporosity while the heat-affected zone (HAZ) showed the presence of microcracks due to the rapid heating and cooling effect.
Other studies have also reported on the formation of various defects in the FZ and HAZ of SC nickel-based superalloys under laser melting are briefly presented: Liang et al. [8] laser remelted a first-generation SC nickel base superalloy SRR 99 with an energy density of 100 J/mm 2 and beam diameter of 2 mm. They found the presence of stray grains in the FZ and attributed their presence to a transition from columnar to equiaxed mode of solidification under the given heat flow conditions. Basak et al. [9] found the formation of fine epitaxial columnar dendrites and equiaxed dendrites in additively manufactured CMSX-4 superalloy processed with a 1 kW ytterbium fiber laser and beam diameter of 40 μm. Nishimoto et al. [10] performed laser remelting experiments on CMSX-4 superalloy with laser power ranging from 150 to 1900 W and scanning speed ranging from 0.01 to 100 mm/s. They found the formation of SC microstructure in the FZ under low heat inputs and poly-crystalline grain structure with stray crystals under high heat inputs. Liu et al. [11] found epitaxial growth of a SC nickel-based superalloy in laser melting using a high-power flat-top laser profile under a longer exposure time.
Mendicoa et al. [12] performed laser cladding experiments on a SC nickelbased superalloy Rene 80 substrate using laser density ranging from 2 to 4800 J/mm 2 with laser spot diameters of 0.2 mm and 2.35 mm using SC Rene 80 powder feed rate ranging from 5 to 30 g/min. The FZ showed a polycrystalline structure with cracks and pores in the deposits under higher heat inputs and SC microstructure under lower heat inputs. Rottwinkel et al. [13] performed laser cladding experiments on SC turbine parts made of CMSX-4 superalloy for crack repair using a solid-state diode laser with an average energy density of 172 J/mm 2 with a fixed laser beam diameter of 0.86 mm, powder feed rate of 3 g/min and z-elevation per layer of 0.15 mm. They found multiple cracks and pores on the surfaces of the clad deposits due to a non-uniform temperature gradient. Liu et al. [14] performed laser powder deposition experiments with a SC nickel-based superalloy Rene N5 powder on a directionally solidified Rene 125 plate by using ytterbium fiber laser with energy densities ranging from 17 to 220 J/mm 2 with a laser beam diameter of 0.6 mm and powder feed rate ranging from 2 to 5 g/min. They found that the columnar dendrites grew epitaxially from the substrate and were stopped by the formation of stray grains. They observed that the height of the epitaxial columnar dendrite zone, as a fraction of the total pool depth, increased with increasing scanning speed, and decreased with an increase in laser power.
Zhou et al. [15] observed that the susceptibility to solidification cracking increased with an increase in the misorientation angle of the grain boundary in a Re-free, SC nickel-based superalloy CSU-B1 fabricated by a 250 W Yb-fiber laser powder deposition at a scanning speed of 8 mm/s. David et al. [16] found that the microstructure of the weld shows the dendrite growth pattern and the presence of stray grains in the canter associated with extensive hot-cracking due to the presence of low melting eutectics in laser welds of a SC nickel-based superalloy, PWA 1480 discs under power ranging from 90 to 240 W, pulse length from 1 to 3 ms, and pulse rate ranging from 20 to 200 s −1 of a Nd:YAG laser with welding speed ranging from 2.1 to 25 mm/s. Korsmik et al. [17] found solidification cracks along the grain boundaries in the clad metal of a second-generation SC nickel-based superalloy ZhS32 due to obstructed shrinkage, under laser energy density of 55 J/mm 2 and a spot diameter of 0.9 mm.
In summary, very few studies have been reported on the evolution of surface defects and their effect on the mechanical strength of the surface of SC nickelbased superalloys by controlling energy densities in laser surface remelting. The mechanisms responsible for the formation of surface defects are very complex and dependent on the initial condition of the alloy and the laser processing parameters. The laser beam diameter is a key parameter in laser surface remelting that significantly influences the incident energy density supplied to the work material and the melting of the surface area. The extent of beam oversize (excess size compared to the size of the beam in focus) depends on the focal position, i.e. the distance of the focus from the surface of the workpiece. The present experimental work investigates the possibility of introducing the controlled laser-induced surface defects and modify the FZ geometry from keyhole to conduction mode penetration, by simply changing the laser focal position in the positive direction, which is easily controllable. The effect of positive focal position on the FZ geometry, microstructure, porosity, cracking, and mechanical strength has been studied, and the related mechanisms have been discussed in this work.

Work Material and Specimen Preparation
The nominal chemical composition of CMSX-4 alloy (by weight) is: Co of 9.6%, Cr of 6.5%, Ta of 6.5%, W of 6.4%, Al of 5.6%, Re of 3.0%, Ti of 1.0%, Mo of 0.6%, and Hf of 0.1% with balance Ni [18]. By using the grain selector technique in combination with directional solidification, the master alloy was cast as a SC turbine blade component with < 001 > crystallographic direction within 15° tolerance from the blade axis through vacuum melting. A solid block was cut by precision WEDM (Wire-Electric Discharge Machining) from the fir-tree portion of the casting, and specimens were gently ground on all surfaces to remove the EDM-related surface defects. The specimens were of length 45 mm, width 16 mm and thickness 13 mm having a flatness of 20 μm, and average surface finish, Ra, in the range of 0.2-0.3 μm on all the surfaces.

Experimental Setup for Laser Surface Remelting
The experimental setup used for the laser treatments consists of four major subsystems which include a high-power CW fiber laser, optical head, shielding gas system, and workpiece positioning system. Figure 1 illustrates the schematic diagram of the experimental setup used for the laser surface remelting experiments on the SC nickel-based superalloy specimens. The CW fiber laser system (model of FL040 from M/s Rofin-SINAR Laser GmbH, Hamburg, Germany), has neodymium ytterbium-doped solid-state laser with an average output power of 4 kW + 1%, a wavelength of 1070 + 10 nm, optical to optical efficiency of around 80%, wall plug efficiency of 30% and a pulse modulation frequency of 5 kHz [19]. The laser system is coupled to the optical head by an optical fiber with a core diameter 200 μm. The longer focal length optics offers a larger depth-of-focus. The optical head has a collimation optics distance of 135 mm and a focusing optics distance of 300 mm. The gas nozzles are fitted to the optical head externally for the supply of the shielding gas. The complete optical head was mounted by a gantry-type system for precise movement in three-linear axes. The workpiece positioning system consisted of a fixed horizontal work table with various work-holding fixtures. The laser power system, the optical head positioning system and the argon shielding system have a computer numerical control (CNC) interface for the selection and execution of various processing parameters.

Laser Beam Characteristics
The laser beam waist diameter is expressed as [20]: where, d 0 is the laser beam waist diameter in μm, d c is the fiber core diameter in μm, M is the magnification factor, it is a ratio of the focal length ( l f ) and collimator lens ( l c ). The value of d 0 that was calculated from the above equation with the installed optics was 440 μm. The laser beam quality is characterized by the beam parameter product (BPP) given by the following equation [21]: where, w o is the beam waist radius in mm and is the beam divergence half-angle in milli-radians. CW fiber laser beams have an excellent beam quality and smaller beam divergence angle. The value of was evaluated through an experimental technique by marking laser spots on a paper using lower power under pulse mode as described in Annexure. The obtained value of as 2.35° was in comparable with the equipment manufacturer's data of BPP of 9 mm.mrad based on w o as 0.22 mm (half value of d 0 = 440 μm, reported above) [19]. The focus depth calculated was 4.5 mm from the Rayleigh length (Z R ), where the Rayleigh diameter ( d R ) increases to √ 2 times of d 0 . A schematic of the CW fiber beam profile along with the focus depth region is shown in Fig. 2.

Experimental Parameters and Procedure
The laser beam diameters corresponding to focal positions ranging from 0 to 5 mm were evaluated from the laser beam profile. The energy density is the combination of the laser power density and laser interaction time. The energy density values under selected focal positions were calculated as the ratio of laser power to the product of scan speed and laser beam diameter [7]. The workpiece was positioned firmly in a fixture. A stand-off distance of 300 mm was maintained between the processing head and the top surface of the workpiece to get zero focal position. High-purity (99.99%) argon with a flow rate of 10 lpm at atmospheric pressure was used as the shielding gas. A laser power of 1000 W and a scan speed of 5.5 mm/s were used for each laser track. The optical head was moved vertically upward to get various positive focal positions. A fixed offset distance of 5 mm between the successive laser tracks was maintained to avoid overlapping of the FZ and HAZ. A schematic of the position of the laser beam d 0 on the surface of the test specimen under various focal positions is shown in Fig. 3a. The laser surface remelting experiments were performed with the parameters mentioned in Table 1 by developing a CNC part program.
The photographs of the laser scans made on the surface of the SC nickel-based superalloy test specimens under the focal positions from 0 to 5 mm are shown in Fig. 3b.

Characterization
Samples from the both start and end zone of the laser scanning were extracted by a precision WEDM cutting perpendicular to the scanning length. These samples were mounted in Bakelite moulds and polished using a series of SiC-coated papers and diamond paste following the standard metallographic procedures to get the cross-sections of the laser scans. Later, the samples were dip etched with Kalling's 2 reagent. Scanning electron microscopy (SEM) coupled with Back Lasers in Manufacturing and Materials Processing (2023) 10:485-521

Fig. 2
Schematic diagram of the CW fiber laser beam profile along with focus depth region Scattered Electron (BSE) imaging was performed to reveal the FZ geometry, microstructures in the FZ, and HAZ. The dimensions of the FZ, pores, dendrites, and crack morphology were measured on the SEM images by using ImageJ opensource software. Qualitative and quantitative elemental analyses were carried out by using an Electron Probe Micro Analyzer (EPMA), EPMA SX-100 model from M/s Cameca, France with 1 μm diameter probe. Hardness measurements in the FZ were carried out by using a Vickers micro-hardness tester, MMTX7 model from M/s MATSUZAVA, Japan with a diamond pyramid indenter under 500 g load and a dwell time of 10 s.

Geometry of the FZ
Low magnification (80 x) SEM micrographs of the cross-section of the laser scans corresponding to various focal positions are presented in Fig. 4. The shape of the cross-section of the laser scans under lower focal positions resembles an ice cream bowl with a long base. At this magnification, the FZ is seen as featureless region with crown, pores, and cracks, whereas the BM beside the FZ, shows a coarse dendritic microstructure. The interface between the FZ and the BM is sharp and the nature of the dendritic microstructure in the BM remains unchanged up to the fusion boundary (FB), which was shown with a white dotted curve, indicating that the HAZ is very small, possibly due to the focussing highly concentrated laser beam and the steep thermal gradient (G), which are typical of laser materials processing. Figure 5a defines important geometrical characteristics of the FZ for the keyhole mode of penetration. Dimensions of the FZ width and the FZ depth along with the aspect (depth to width) ratio in changing the focal position is shown in Fig. 5b. The FZ under keyhole mode has a moderate level of crown formation. It is known that during keyhole formation, due to the large G existing in the molten pool, the associated recoil pressure and the Marangoni effect force the melt to flow from the bottom to the top, leading to crown formation [22]. The width of the FZ is slightly larger than the depth for all the focal positions and both reduce with increasing focal position with the latter decreasing more steeply beyond the focal position of 3 mm. As a result, the aspect ratio of the FZ decreases from 0.8 to 0.3 with an increase in the focal position. As the focal position increases, the effective laser energy density decreases, and FZ depth reduces and weakens the melt flow described above, eventually eliminating the crown as observed, e.g., at 5 mm focal position in the present case. Dimensions of the keyhole (KH) geometry such as KH width, KH depth, and base radius with changing focal position are shown in Fig. 5c. Both KH width  and KH depth also decrease with increasing the focal position. The shape of the FB at the base is half circular and the radius of the FB also decreases linearly with increases in focal position. This could be due to the multi-mode Gaussian beam profile of a CW fiber laser employed in the present work. The aspect ratio of keyhole geometry consists of higher values ranging from 1.3 to 1.2 with increasing the focal position. As a whole, the cross-sectional area of FZ also decreases with increasing focal position due to a decrease in the energy density (Fig. 5d).
When the focal positions are up to 3 mm (with energy densities ranging from 410 to 260 J/mm2), the keyhole mode of penetration occurs, leading to a decrease in KH geometry (long base) and overall FZ size. On the other hand, when the focal position is 5 mm (with an energy density of 235 J/mm2), the conduction mode of penetration occurs, resulting in a short base geometry of the FZ. The depth of FZ as a function of estimated incident energy density along with the three modes of penetrations under observed focal positions are shown in Fig. 6.
These observations are commensurate with the known facts that the laser energy density primarily controls the penetration depth of the melt pool during laser processing [23,24]. Assuncao et al. [25] found that CW laser weld pools have a surface tension gradient and the associated melt movement, which are responsible for the transition mode region, whereas the pulse wave laser welds have only the surface tension gradient and but no melt pool movement, which resists the formation of transition mode of penetration. Fabbro et al. [26] observed that in the keyhole mode, the vapor pressure recoil displaces the molten material and thereby increases the penetration depth.

Microstructure
High magnification (2000 x) SEM micrographs within the FZ near the central location (un-cracked region) and at the FB corresponding to various focal positions are shown in Fig. 7a to f. The FZ has a fine dendritic microstructure, with the dendrites seen in various orientations. In some cases, their primary arms lie parallel to the section, while in others they are oriented perpendicular to the section showing the star-shaped (or cruciform) dendritic cross-sections. It is possible that some of the cruciform-shaped dendrites are, in fact, remelted parts of secondary arms, but it was not possible to confirm this. In yet other cases, dendrite primary arms are not seen and only the secondary arms are seen as islands or in nodular shapes, possibly because the section plane did not pass through the primary arm. The number of dendrites appearing in the FZ increases and, consequently, the primary dendrite arm distance decreases from 16 to 8 μm with an increase in the focal position from 0 to 5 mm. Increasing the focal position reduces the laser energy density which increases the cooling rate and the nucleation rate of dendrites during solidification. Solidification is primarily epitaxial along the FB where a featureless layer of solid, referred to as the 'beach region' [27] in a recent publication, is observed. In some cases, columnar dendrites seem to originate from this beach region. The thickness of the beach region decreases from 12 to 3 μm with an increase in the focus position. The beach region is associated with 1 3  (continued) a partially melted region formed near the FB, which does not fully mix with the bulk of the FZ. This partially melting region could be formed by the dissolution of the low melting eutectic compound present in the BM nearer to the FB during the solidification. With increasing focal position, the energy density decreases, reducing the thickness of the partially melted zone and, hence, that of the beach region. The FB consists of microcracks and dense micropores. The presence of micropores in the inter-dendrite region led to the development of microcracks along the dendrites in the FZ. A schematic diagram of the microstructural changes in the keyhole mode and conduction mode in the FZ of a SC nickel-based superalloy under laser surface remelting is shown in Fig. 8. It is observed that the base region of the FZ in key hole mode has much finer dendrites compared to those in conduction mode. This may be attributed to the higher cooling rates in the smaller area of the base region.

Porosity
High magnification SEM micrographs of the pores in the FZ are shown in Fig. 9. It is found that the FZ has microporosity with a size of the order of 2-10 μm. Apart from this, coarse porosity observed is in the size range of 60-80 μm, but fewer in number is seen somewhat near the FB. While the coarser porosity of size around 200 μm is observed in the middle of the FZ under the keyhole mode, in contrast, no such pores are observed in the conductive mode. Norris et al. [28] found that the keyhole instabilities and the associated melt flow result in the formation of porosity. Katayama et al. [29] found that a deep keyhole collapses under low welding speeds. When the laser beam is irradiated on the collapsed wall of the keyhole, downward melt flow occurs by the recoil pressure of evaporation from the collapsed wall due to which large bubbles form in the melt, which give rise to pores after solidification. If the energy density increases due to lower scan speeds and/or higher laser power, it results in the boiling of low melting eutectics and entrapment of vapour in the keyhole region which is ultimately responsible for the formation of large pores.   Figure 7 shows that certain microcracks exist in the FZ near the FB, mostly the microcracks are in parallel and perpendicular directions to the FB. The magnified images of microcracks in keyhole mode, as well as conduction mode are shown in Fig. 10. The microcracks in keyhole mode have wider crack width (1 to 2 µm) than the conduction mode. But, the number of microcracks in conduction mode was higher in number. The shallow and dense microcracks in conduction mode could be due to higher cooling rates under lower energy density. The microcracks were also developed in the brittle eutectics nearer to the FB due to the thermal stresses of the FZ.

Solidification Cracks
The magnified views of the solidification cracks in the FZ of keyhole mode and conduction mode are seen in the SEM micrographs for the focal positions of 1 mm, 3 mm, 4 mm, and 5 mm respectively in Fig. 11. The vertical cracks are seen to be near the centreline of FZ, contain fine globular solid constituents, which are possibly fragments of dendrites. Usually, these fragments originate by re-melting and convective transport of secondary dendrite arms.
The SEM micrographs of the solidification cracks produced in FZ under the keyhole and the conductive modes are shown in Fig. 12a and b, respectively. It can be seen that the FZ of keyhole mode has vertical as well as radial cracks in the base region while only vertical cracks nearer to the surface are seen in the conduction mode.  The crack formation mechanisms in the keyhole and conduction modes are illustrated in Fig. 13. The vertical cracks are believed to form due to thermal stresses generated due to constraints on the thermal contraction of dendrites during solidification along the principal direction of heat extraction which is parallel to the melt pool width. Weld centreline cracks are commonly described in welding literature [30]. Since the base region has circular FB, solidification proceeds from the circular boundary towards the centre (within the base region), and the thermal stresses associated with the restrained contraction are expected to be circumferential, resulting in radial cracks.
The crack dimensions such as crack width, the total crack length (TCL), and TCL per unit area of the FZ were evaluated from both the start and end zones of the laser scan. The crack dimensions as a function of focal position are shown in Fig. 14. It can be seen that the crack width has a higher value with a large error bar under the focus position on the surface (0 mm). The TCL and TCL per unit area decrease linearly with an increase in focal position from 0 to 4 mm (keyhole mode to transition mode). The value of TCL and TCL per unit area decreases drastically under the focal position of 5 mm (conduction mode). The severity of cracking increased with increasing laser energy density which, in turn, promoted the keyhole mode of interaction. The keyhole mode is expected to generate a more complex thermal stress distribution, resulting in more severe cracking.

Elemental Segregation
Elemental microsegregation studies were performed on the cross-section of the laser scans at the dendrite core (DC) and the inter-dendrite (ID) regions of the cruciformshaped dendrites adjacent to the crack region in the FZ for analysing the local chemistry. The corresponding probing locations in the BSE image nearer to the surface and keyhole region in the FZ from the specimen laser treated with a focal position of 1 mm are shown in Fig. 15. The segregation ratio of the alloying elements was   Fig. 15c. The DC region shows higher values of dense refractory elements (Re, W, Co, and Mo) and lower values of γ'-forming elements (Al, Ti, and Ta), and the opposite trend is seen in the ID region. The scatter in the segregation ratio is seen to be higher in the ID region than in the DC region. Overall, these observations are in conformation to the known partitioning behaviour of  various elements during non-equilibrium solidification, especially so for SC nickelbase superalloys containing slow diffusing refractory elements like Re, W, and Mo [31].
Next, it is sought to compare the elemental partitioning between the DC and the ID regions, which results due to limited elemental diffusion in the solid state for the BM and the laser FZ. The degree of partitioning is commonly expressed by the partition coefficient (k o ) as a ratio of the elemental concentration in the DC to that in the ID region. Elemental concentrations were measured by EPMA and the k o values corresponding to various alloying elements were calculated (Fig. 16). It is observed that the k o value for all the elements except the Ti element is lesser than the BM, indicating a lower level of microsegregation. For example, for the slow diffusing refractory elements which have a positive value of the partitioning coefficient, there was a 19% reduction in k o for Re, 16% for Mo, and 6% for W in the FZ when  Further, elemental distribution was evaluated near the solidification cracking and microcracking regions in FZ and eutectic regions in the BM. Figure 17a shows the complete cross-section of a keyhole-type scan made using a focal position of 3 mm, marking the crack edge (CE1) at the vertical crack and the CE2 at the radial crack and Fig. 17b shows the enlarged view of selected region at FB, marking the CE3 at microcracks and various eutectic regions in the BM. The segregation ratios of the alloying elements are shown in Fig. 17c. It is observed that the segregation ratios at CE1 and CE2, which are in the bulk of the FZ, are close to one, whereas CE3, which is located near the FB, is enriched in Ti, Ta, Al, and significant segregation near the cracks in the central region of the FZ, which suggests that the solidification cracks were driven by high thermal stress generated near the melt pool centreline. The eutectic in the HAZ is enriched in Ta and Ti and depleted in refractory elements such as Re, W, and Mo. The segregation study was further extended to the stray solids which were observed within the solidification cracks in the FZ. Figure 18a and c show the cross-section of the keyhole scan made using focal positions of 1 mm and 3 mm respectively, marking the selected area for the study. Figure 18b and d show the enlarged views of the solidification cracks in the FZ, marking the probing locations for the elemental analysis. Figure 18e shows the segregation ratios of the alloying elements of the stray solids. It is observed that the stray solids within the crack are enriched in Re, Mo, and W, and deficient in Al and Ti, which suggests that these are fragments of dendrite arms re-melted at their roots and transported into the crack region by convection. Figure 19 shows the cross-section of the FZ of a keyhole-type laser scan where a vertical crack extending from the surface to the interior and several radial cracks in the base region are seen. An area of the crack near the surface, as shown in Fig. 19a, was selected for elemental analysis. Figure 19b shows the BSE image from the selected area and Fig. 19c to k show the corresponding Fig. 19 Elemental maps of a solidification crack (a) cross-sectional view of the FZ of a keyhole mode laser scan made using a focal position of 1 mm, marking a small area of the solidification crack selected for X-ray maps studies, (b) BSE image of the selected area, and (c) to (k) X-ray elemental maps of various elements as labeled 1 3 X-ray maps for various alloying elements. There is an enrichment of refractory elements W and Re, along with Ni, Co, and Cr in the stray solids, while the dark phase is enriched in Al, Ti, and Ta. The X-ray signal for Mo is weak making it difficult to conclude its segregation.
The elemental analysis in the area at the center of the radial crack in the base region of the keyhole type FZ produced under a 1 mm focal position showed that the segregation levels of various alloying elements are at slightly reduced levels of intensity (figure not included). The reduced segregation levels in the base region compared to the top region of the FZ could be due to faster cooling rates prevailing in the base region than in the top zone. Overall, the presence of micro-segregation (resulting in lowering of the non-equilibrium solidus temperature) coupled with thermal stresses are considered to be the reasons for cracking in the base region of the keyhole mode FZ.

Microhardness
Bulk hardness in the FZ was not measured, although it is expected to be lower than the BM due to the presence of cracks and pores. It was sought to know if the microstructural changes in the FZ and near the FB had any effect on the sound portions of the FZ. Hence, Vickers micro-hardness indentation tests were performed beneath the surface up to 3.6 mm covering the FZ and beyond FZ as per ASTM standard E384-17 for the laser scans made with various focal positions [32]. The locations of hardness measurements beneath the surface are shown in the form of indentations in Fig. 20a to f which correspond to various focal positions under which the laser scans were made.
There was a scatter in the hardness values within the FZ to the extent of ± 10% from the average value as shown in Fig. 21a and b plots the average hardness values along with error bars within the FZ and below the FZ (i.e. in the BM) for laser scans made using the focal positions from 0 to 5 mm. The figure shows that the average hardness in FZ is about 5 to 7%) lesser than below FZ for all focal positions except at 0 mm. The close matching FZ with the below FZ at 0 mm could be due to either a higher value of microhardness in the FZ or a lower value in the below FZ. Thus, the laser surface melting has a clear advantage of reducing the hardness even outside the defects within the FZ. The drop in the hardness can be attributed partly to the dissolution of the strengthening ϒ' precipitates and to the presence of defects such as micro-pores and microcracks in the vicinity. The hardness of the FZ under conduction mode is found to be marginally lower than in keyhole mode (Fig. 21c), it could be due to an increase in defect formation (micro-cracks and micro-porosity) under higher cooling rates which are likely to be associated with lower energy density.

Summary and Conclusions
Laser surface remelting experiments were performed on a commercial second-generation SC nickel-based superalloy CMSX-4 with a range of focal positions from 5 to 0 mm at a constant laser power of 1000 W and a scan speed of 5.5 mm/s. The following are the major findings of the experimental investigation.
• The use of energy densities above 260 J/mm 2 , corresponding to focal positions less than 3 mm, generated the keyhole mode, whereas energy densities lower than 235 J/mm 2 corresponding to the focal position of above 4 mm resulted in conduction mode penetration. The transition mode prevailed for the intermediate energy density. The dimensions of the FZ such as the width, depth, aspect ratio, and cross-sectional area decreased with an increase in focal position due to a decrease in energy density. • The FZ is characterized by the presence of pores, cracks, and assorted fine dendrites, a substantial absence of the eutectic and the γ′phase, and lower levels of elemental segregation ratios. The central region shows cruciformshaped dendrites, appearing so due to the direction of dendrite growth at the centreline being perpendicular to the plane of viewing, while the FB region shows columnar dendrites. The scale of the dendritic microstructure becomes finer with an increase in focal position due to enhanced nucleation under higher cooling rates. The planar dendrite is observed near the FB, the thickness of which reduces with an increase in the focal position. Coarse pores formed in the base region under keyhole penetration and in contrast, no such pores are observed in the conductive mode. • The FB is sharper and has no visible HAZ, mainly due to the focussing highly concentrated laser beam and the steep thermal gradient (G) expected during laser material processing. The eutectic in the HAZ is enriched in Ta and Ti and depleted in refractory elements such as Re, W, and Mo. • There is no significant segregation near the crack edges in the central region of the FZ, which suggests that the cracks were formed after solidification by the high thermal stresses. The cracks under the conductive mode were mostly vertical centreline cracks, while under keyhole penetration, radial cracks formed in the base region, in addition to the centreline vertical cracks. Fine solid globules are observed inside the solidification cracking. The EMPA analysis confirms that these solid globules are a mixture of highly dense refractory elements and fragments of dendrites. The formation of cracks is aided by a lack of feeding solidification shrinkage near the FB. Geometric parameters like the crack width, total crack length, and crack areal density decreased with an increase in the focal position. • Microcracks were developed near the prevailing eutectic phases at the FB. The keyhole mode consists of intense microcracks, whereas the conduction mode consists of denser microcracks due to higher cooling rates under lower energy density. The significant level of segregation in the microcrack regions suggests that the incipient solidus temperature must have reduced locally, increasing the solidification temperature range, and making the region vulnerable to solidification cracking.
Thus, The present experimental shows that it is possible to modify the FZ geometry from keyhole to conduction mode penetration, by simply changing the laser focal position in the positive direction, which is easily controllable.
The FZ hardness in the crack-free regions was lower than the BM, which may be partly due to the dissolution of γ′ phases and the formation of multiple laserinduced surface defects described above under laser remelting. There is an expectation that the bulk hardness of the FZ (which includes the defect) would be even lower which will make the laser-remelted surface more amenable to machining.

Results and discussion
The optical images of laser spots on a graph sheet under focal positions in ascending order from 0 to 8.5 mm with laser power of 400 W and time duration of 100 µs are shown in Fig. 23.
The optical images of laser spots on a graph paper under AFD values in descending order from 8.5 to 0 mm with laser power of 400 W and time duration of 100 µs are shown in Fig. 24.
The results of laser spot diameter on a graph paper with ascending and descending orders of focal positions and the comparative results of laser beam diameter and average laser spot diameter are shown Fig. 25. The below equation is best fitted for the average spot diameter by the straight-line geometrical equation: where, x is the focal position in mm and y is the spot radius in mm. It can be seen that the slop of the laser beam diameters evaluated from the theoretical values by using the manufacturer catalogue value closely matches with the slope of the laser spot evaluated through the graph sheet method. The achieved slop was converted into the degrees through below equation to get divergence half-angle ( ) Therefore, the evolved divergence half-angle of the CW fiber laser beam is 2.35°.