Modern fiber laser beam welding of the newly-designed precipitation-strengthened nickel-base superalloys

Modern fiber laser beam welding of the newly-designed precipitation-strengthened nickel-base superalloys

Optics & Laser Technology 57 (2014) 12–20 Contents lists available at ScienceDirect Optics & Laser Technology journal homepage: www.elsevier.com/loc...

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Optics & Laser Technology 57 (2014) 12–20

Contents lists available at ScienceDirect

Optics & Laser Technology journal homepage: www.elsevier.com/locate/optlastec

Modern fiber laser beam welding of the newly-designed precipitation-strengthened nickel-base superalloys Homam Naffakh Moosavy a,n, Mohammad-Reza Aboutalebi a, Seyed Hossein Seyedein a, Massoud Goodarzi a, Meisam Khodabakhshi b, Carlo Mapelli c, Silvia Barella c a

School of Metallurgy and Materials Engineering, Iran University of Science and Technology (IUST), 16846-13114 Tehran, Iran MAPNA Turbine Engineering and Manufacturing Company (TUGA), 15875-5643 Fardis, Karaj, Iran c Dipartimento di Meccanica, Politecnico di Milano, Via La Massa 34, 20156 Milan, Italy b

art ic l e i nf o

a b s t r a c t

Article history: Received 30 April 2013 Received in revised form 6 September 2013 Accepted 21 September 2013 Available online 15 October 2013

In the present research, the modern fiber laser beam welding of newly-designed precipitationstrengthened nickel-base superalloys using various welding parameters in constant heat input has been investigated. Five nickel-base superalloys with various Ti and Nb contents were designed and produced by Vacuum Induction Melting furnace. The fiber laser beam welding operations were performed in constant heat input (100 J mm  2) and different welding powers (400 and 1000 W) and velocities (40 and 100 mm s  1) using 6-axis anthropomorphic robot. The macro- and micro-structural features, weld defects, chemical composition and mechanical property of 3.2 mm weldments were assessed utilizing optical and scanning electron microscopes equipped with EDS analysis and microhardness tester. The results showed that welding with higher powers can create higher penetration-to-width ratios. The porosity formation was increased when the welding powers and velocities were increased. None of the welds displayed hot solidification and liquation cracks in 400 and 1000 W welding powers, but liquation phenomenon was observed in all the heat-affected zones. With increasing the Nb content of the superalloys the liquation length was increased. The changing of the welding power and velocity did not alter the hardness property of the welds. The hardness of welds decreased when the Ti content declined in the composition of superalloys. Finally, the 400 and 1000 W fiber laser powers with velocity of 40 and 100 m ms  1 have been offered for hot crack-free welding of the thin sheet of newly-designed precipitation-strengthened nickel-base superalloys. & 2013 Elsevier Ltd. All rights reserved.

Keywords: Modern fiber laser beam welding Precipitation-strengthened nickel-base superalloys Weld defects

1. Introduction Precipitation-strengthened nickel-base superalloys are extensively utilized in the important industrial applications such as land-based and aero gas turbines due to their excellent high temperature strength. These alloys mainly obtain their high temperature strength due to formation of dispersive γ′-Ni3(Ti,Al) and γ″-Ni3Nb intermetallic particles in the austenite phase [1]. High content of titanium and niobium in the composition of nickel-base superalloys not only causes the γ′ and γ″ strengthening phases to form, but also can lead to formation of various types of other phases and secondary precipitations in the microstructure. Carbide precipitations such as TiC, NbC, and (Ti,Nb)C, Laves eutectic structures such as γ-Ni3Nb and γ-Ni3(Ti,Al) and other eutectic structures like γ-NbC, or δ-Ni3Nb intermetallic compounds are examples of these secondary phases [2–5]. Secondary

n

Corresponding author. Tel.: þ 98 21 77240540x2816; fax: þ 98 21 77240480. E-mail address: [email protected] (H. Naffakh Moosavy).

0030-3992/$ - see front matter & 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.optlastec.2013.09.030

phases and structures can strongly deteriorate the weldability of superalloys. TiC, NbC and (Ti,Nb)C precipitations melt when they react with the austenite phase and reach to the eutectic composition. These molten regions can facilitate the liquation cracking in the heat-affected zone of the superalloy by constitutional liquation mechanism. Austenite–Laves and austenite–carbide liquid eutectic structures solidify at lower temperatures in comparison to the austenite phase solidification temperature [6–8]. Thus, these compounds increase the solidification temperature range of the superalloys and sensitize them to the hot solidification cracking. The solidification temperature range (ΔT) is measured by the ΔT¼ Tl  Ts, in which, Tl is the liquidus, and Ts is the solidus temperatures. When the final liquid of the weld metal solidifies as the above eutectics, the liquidus temperature of weld decreases, because the temperature of eutectic formation is significantly lower than formation temperature of austenite single phase. As a result, the solidification temperature range (ΔT) of the weld containing the above eutectics increases. When these phases are present in the heat-affected zone (HAZ) of the welds, the liquation cracking occurs by remained eutectic mechanism melted during

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Table 1 Chemical composition of the newly-designed superalloys. Alloy designation

Alloy Alloy Alloy Alloy Alloy

1 2 3 4 5

Composition of elements (wt%) Ni

Cr

Co

Fe

Mo

Al

Ti

Nb

52.7 52.4 52 51.8 51.4

19 19 19 19 19

14.8 11.1 7.4 3.7 –

– 4.6 9.4 13.9 18.6

3.5 3.5 3.5 3.5 3.5

3 2.4 1.8 1.2 0.6

7 5.5 4 2.5 1.1

– 1.5 2.9 4.4 5.8

the welding. Moreover, this is reported that the γ′ and γ″ strengthening precipitations and δ phase cause liquation cracking in the HAZ by constitutional liquation mechanism. Therefore, the precipitation-strengthened nickel-base superalloys are severely susceptible to the hot cracking, and significant efforts are carried out to design new superalloys by optimizing the alloy chemical composition in order to achieve improved weldability [9–12]. On the other hand, considerable researches have been performed to substitute the conventional welding processes by the newer welding methods for welding the nickel-base superalloys [13–17]. Usual welding processes such as Shielded Metal Arc Welding (SMAW), Gas Metal Arc Welding (GMAW), and Gas Tungsten Arc Welding (GTAW) can impose significant damages to the nickel-base superalloys weldability due to low heat concentration; low heat transfer efficiency and wide weld metal and heat-affected zones. Thus, the welding processes with high energy concentration and deep penetration-to-width ratio are strongly spreading for joining the susceptible nickel-base superalloys [18–20]. Beam welding processes such as Electron Beam Welding (EBW) and Laser Beam Welding (LBW) are examples for the new welding techniques with deep penetration. Appropriate penetration, low heat input, small weld and heat-affected zone, improved efficiency, clean welding, welding speed and thin sheet welding are the most important advantages of the beam welding processes. The researchers correlated the laser welding heat inputs to the size of heat-affected zone, the liquation cracking and solidification cracking in the Inconel 625 and Inconel 718 nickel-base superalloys such. They found that in the lower heat inputs, smaller heat-affected zones and cracks are obtained [21–23]. The laser beam welding techniques such as CO2laser and Nd:YAG laser have found more applications for nickel-base superalloys welding than electron beam welding [24–26]. Nowadays, with innovation of the Fiber Laser Beam (FLB), modern welding of the nickel-base superalloys with the fiber laser beam has been provided. However, the Fiber Laser Beam Welding (FLBW) of the precipitation-strengthened nickel-base superalloys has not experimentally been studied. So, the purpose of the present research is systematic investigation of the welding and weldability of the newlydesigned superalloys welding using the modern fiber laser beam welding. In this organized research, the fiber laser beam welding of newly-designed precipitation-strengthened nickel-base superalloys has been fulfilled utilizing various welding powers and welding velocities in constant heat input. The geometrical evaluations of the welds (i.e. penetration-to-width ratio), microstructural aspects, weld defect assessments, weldability evaluation and mechanical investigation of fiber laser beam-welded newly-designed superalloys have been carried out in detail.

2. Experimental procedure One of the current research aims is to investigate the effect of Nb,Ti variations on the microstructure and weldability of obtained fiber laser-welded superalloys. The precipitation-strengthened superalloys are specially hardened by the gamma-prime and gamma-double

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prime intermetallics which are enriched in the above mentioned elements. The other elements such as Ni, Co, Fe, Cr and Al form the austenite matrix, whose variations in the composition of superalloys have minor effects on the microstructure and weldability. Therefore, five different nickel-base superalloys with pre-specified composition have been designed. A vacuum induction melting (VIM) furnace with the vacuum of 10  4 bar, the frequency of 3.6 KHz and maximum power of 60 KW was employed for melting and adjusting the composition of the superalloys. Pure alumina crucibles with volume of 1000 ml, and steel die with dimensions of 150 mm  100 mm  50 mm were used for pouring the resultant molten alloys. The temperature of 1450 1C was selected for full melting of alloys, and the ingots were cooled at the ambient temperature. The chemical composition of newly-designed superalloys is shown in Table 1. The Ti content decreases and the Nb content increases from Alloy 1 to Alloy 5. Before the welding, the strips with dimensions of 150 mm  25 mm  3.2 mm were extracted from the newly-designed as-cast superalloy ingots by wire cutting. Then, the samples were deoxidized and degreased by a stainless steel brush and acetone solution. The Fiber Laser Beam Welding equipment was utilized without adding the filler metal (i.e. autogeneous welding) to produce samples containing bead-on-plate welds. The Fiber Laser Beam Welding (FLBW) set up was composed of an IPG YLR-1000 fiber laser source with maximum power of 1000 W, frequency of 5 kHz, and 1070 nm wavelength; a HIGHYAG BIMO welding head with 200 mm focus lens diameter and 100 mm collector; and an ABB IRB 2400 6-axis anthropomorphic robot as a positioning machine. The essential parts of FLBW setup are shown in Fig. 1. Two different laser powers and welding velocity values were used in constant input energy applied to the weld pool. The details of FLBW parameters are provided in Table 2. The input heat energy per unit square millimeter (i.e. input heat surface density) was calculated on the basis of the following equations:   J P ðJ=sÞ E ð1Þ ¼ V ðmm=sÞ  d ðmmÞ mm2 For P ¼ 400 W- E ¼

400 ðJ=sÞ ¼ 100 40 ðmm=sÞ  0:1 ðmmÞ

For P ¼ 1000 W- E ¼



J mm2

1000 ðJ=sÞ ¼ 100 100 ðmm=sÞ  0:1 ðmmÞ





J mm2

ð2Þ  ð3Þ

where P is fiber laser beam power, E is input heat energy, V is speed of welding, and d is the focused laser beam diameter on the surface. After the welding operations, abrasive cutting and cleaning, the specimens were prepared according to the standard metallographic procedures. At least 3 different perpendicular cross-sections from each weld are used for the microscopic and mechanical evaluations, and the average values are reported as the results. For revealing the weldment microstructures, four different etching solutions were used. The type, chemical composition and the corrosivity of the etchants have been presented in Table 3. The identification of the microstructural characteristics was carried out by a Leitz Wetzlar Aristomet optical microscope equipped with Nikon ACT Version 2.70 software. Further microstructural investigations were performed by a Zeiss EVO 50XVP scanning electron microscope (SEM) with accelerating voltage of 30 kV equipped with an Oxford Instrument 7060 energy dispersive X-ray spectrometer (EDS) for spot, line and map weight analysis. For measurements of mechanical properties i.e. hardness, a Vickers Microhardness Tester of Future-Tech Corporation FM-700 model, with indentation weight of 1000 gf and indentation time of 15 s was utilized. The Vickers microhardness tests were fulfilled in

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Fig. 1. The parts of Fiber Laser Beam Welding (FLBW) set up: (a) IPG YLR-1000 fiber laser source with maximum power of 1000 W, frequency of 5 kHz, and 1070 nm wavelength, (b) the HIGHYAG BIMO welding head with 200 mm focus lens diameter and 100 mm collector, and (c) the ABB IRB 2400 6-axis anthropomorphic robot as a positioning machine.

Table 2 Fiber Laser Beam Welding parameters. Laser power (W)

Argon shielding gas flow rate (l min  1)

Welding velocity (mm s  1)

Focused beam diameter on the surface (mm)

Energy input (J mm  2)

400 1000

15 15

40 100

0.1 0.1

100 100

Table 3 The type, chemical composition and corrosivity of used etchants.

1 2 3 4

The etchant type

The etchant chemical composition

Corrosivity

Glyceregia Marble Modified marble Aqua Regia

30 ml H2Oþ 60 ml HClþ 20 ml HNO3 5 g CuSO4 þ100 ml H2Oþ 100 ml HCl 25 g CuSO4 þ50 ml H2Oþ 150 ml HCl 20 ml HNO3 þ60 ml HCl

Weak Moderate Strong Very strong

the base metals, weld metals and heat-affected zones of the weldments.

3. Results and discussion 3.1. Geometric and macroscopic evaluation of the welds Figs. 2 and 3 show the macrostructures of superalloy welds produced by fiber laser beam welding. These welds have been obtained with 400 and 1000 W powers, and 40 and 100 mm/s velocities, respectively. As seen, the weld region is formed in all newly-designed superalloys. The macrostructure of welds is used

for geometric measurements. Table 4 exhibits the total geometrical characteristics of the welds in the laser-welded samples. All the welds produced with 1000 W laser power indicate key-hole penetration mode. However, only welds of Alloy 2 and Alloy 3 show key-hole mode at 400 W laser power. These results show that key-hole mode can be achieved for the welds by increasing the laser power, even when the laser heat input is constant. The penetration-to-width ratio of welds is calculated and reported in Table 4. The penetration-to-width ratio of the welds versus the superalloy designation has been plotted in Fig. 4. According to Fig. 4, the highest and the least ratios are obtained for Alloy 2 and Alloy 1, respectively. Although, the heat input is constant, the penetration-to-width ratio of the welds for 1000 W laser power is more than that for 400 W laser power applied. Thus, the penetration-to-width ratio of laser-welded superalloys improves when the laser power increases, even in fixed heat input. The results show that Alloy 1 which contains the highest titanium content has the least penetration-to-width ratio. This result can be attributed to the high melting temperature and alloy hardness. Based on Fig. 4, although the penetration-to-width ratio is not the same for the newly-designed superalloy welds, the governing trend is kept relatively fixed for the various laser welding powers. In other words, the penetration-to-width ratio

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Fig. 2. Macrostructure of fiber laser beam-welded newly-designed superalloys in P¼ 400 W, V ¼40 mm s  1: (a) Alloy 1, (b) Alloy 2, (c) Alloy 3, (d) Alloy 4, and (e) Alloy 5.

is remarkably a function of chemical composition and laser power. On the basis of the current results, it can be said that the penetration-to-width ratio of the welds improves when the laser power and the welding velocity increase. This conclusion has been achieved with constant welding heat input of 100 J/mm2.

3.2. Evaluation of the defects and microstructure of the welds The measurements of weld defects in the fiber laser beamwelded superalloys are summarized in Table 5. Weld porosities are observed in all the superalloys welded at 1000 W laser power and 100 mm s  1 velocity (Fig. 3). None of the welds (with exception of Alloy 3) produced at 400 W laser power and 40 mm s  1 display porosities (Fig. 2). The results of Table 5 indicate that formation of weld porosities is directly a function of the laser welding power and speed. The porosity formation is principally a function of heat input. But, in the current work the input energy is constant. Under this condition, the laser welding speed and laser power control the porosity formation. With increasing laser power and welding speed the porosity formation, number of porosities, and the porosity diameter have been increased at constant heat input. High laser power and high welding speed can increase the turbulence in the weld pool and facilitate the entrapment of the argon bubbles in the weld dendritic structure. On the other hand, the number of porosities and the maximum diameter of weld porosities with 1000 W laser power welding versus the superalloy designation are plotted in Fig. 5. The weld of Alloy 3 displays the maximum number and the biggest porosities, and weld of Alloy 5 exhibits the least number and the smallest porosities.

The weld microstructure of fiber laser beam-welded superalloys is shown in Fig. 6. The precise investigation of weld metals shows no evidence of solidification crack formation, either with the 400 W or in 1000 W powers. Since these superalloys are severely sensitive to the solidification cracking, using a welding process and welding power which present welds free of solidification cracks is very valuable. As observed in Fig. 6, the Alloy 1 and Alloy 2 welds exhibit fully-dendritic and continuous microstructures at both 400 and 1000 W laser power. Titanium in the Alloy 1 weld, and niobium in the Alloy 5 weld cause significant microsegregation and full dendritic boundary coalescence. The lighter regions are the microsegregation boundaries and the darker regions are the dendritic cores. Microsegregation partially occurred in the Alloy 2 weld and the dendritic boundaries are thoroughly discontinuous. The weld of Alloys 3 and 4 indicate stronger microsegregation in comparison to weld 2. 3.3. Evaluation of the defects and microstructure of the heat-affected zones Figs. 7–11 display the heat-affected zone microstructure of the laser-welded superalloys. The HAZ is formed in the all superalloy weldments, but no liquation cracking is observed. Nevertheless, all the fiber laser-welded superalloys indicate the liquated regions in the partially melted zone (PMZ) at 400 and 1000 W laser power. The superalloy welds which are made by higher heat inputs are more susceptible to the liquation cracking due to smaller temperature gradient and more extensive stress distribution. However, the welding parameters here did not lead to the liquation cracking. It reveals

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Fig. 3. Macrostructure of fiber laser beam-welded newly-designed superalloys in P¼ 1000 W, V ¼ 100 mm s  1: (a) Alloy 1, (b) Alloy 2, (c) Alloy 3, (d) Alloy 4, and (e) Alloy 5.

Table 4 Weld penetration mode, weld penetration, weld width and penetration-to-width ratio for the fiber laser welded newly-designed superalloys. Laser power (W) 400

1000

Alloy Weld designation penetration mode Alloy Alloy Alloy Alloy Alloy Alloy Alloy Alloy Alloy Alloy

1 2 3 4 5 1 2 3 4 5

– – Key-hole – – Key-hole Key-hole Key-hole Key-hole Key-hole

Weld penetration (μm)

Weld width (μm)

Penetrationto-width ratio

230 700 736 713 586 1057 1161 1025 782 1220

782 552 678 828 586 874 644 736 667 701

0.29 1.27 1.09 0.86 1 1.21 1.80 1.39 1.17 1.74 Fig. 4. The penetration-to-width ratio of the newly-designed superalloy welds, welding with 400 and 1000 W laser powers.

that the welding parameters used could not produce sufficient strain for the cracking. Of course, at the higher heat inputs or in other welding processes such as GTAW and GMAW, the liquation cracking is unavoidable. Some of the molten dendritic boundaries in the PMZ of laser-welded Alloy 1 are enriched in Ti and some other regions are enriched in Cr and Mo. The EDS analysis confirms the claim above (Fig. 7(c) and (d)). These eutectic phases have been liquated by the mechanism of remaining eutectics melting and formed the liquated regions in the heat-affected zone. Liquation by the remaining eutectics mechanism has been explained in the previous works [1,12,16]. The partially melted zone of laser-welded Alloy 2 shows the liquation of Tiand Nb-rich carbides by the constitutional mechanism, and liquation of the Nb- and Mo-rich eutectic structures by melting the remaining

eutectics mechanism (Fig. 8(c) and (d)). The constitutional mechanism in the HAZ of nickel-base superalloys is presented in the investigations [9,11,22]. Alloy 3 indicates the liquation mechanism in the HAZ similar to that of Alloy 2 (Fig. 9(c and d). The melting of Nb- and Mo-rich Laves eutectics occurs in the interdendritic regions of Alloy 4 HAZ (Fig. 10(c)). The HAZ of Alloy 5 displays the melting of the Nb-rich δ intermetallics by constitutional mechanism (Fig. 11(c)). The acicular Ni3Nb precipitations can be seen in the vicinity of the liquated regions. These findings are in agreement with the previous works [21–23].

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Table 5 Summary of hot cracks in the weld and heat-affected zone (HAZ), and the weld porosity for fiber laser welded newly-designed superalloys. Laser power (W)

400

1000

Alloy Number of designation porosities in the weld metal

Max. porosities diameter (μm)

Weld metal solidification cracks

Max. HAZ liquation length (μm)

Alloy Alloy Alloy Alloy Alloy Alloy Alloy Alloy Alloy Alloy

– – 74 – – 114 74 134 66 20

– – – – – – – – – –

20 31 43 47 61 20 31 43 47 61

1 2 3 4 5 1 2 3 4 5

0 0 1 0 0 3 2 3 3 1

Fig. 5. Number of porosities and maximum diameter of weld porosities versus superalloy designation in 1000 W laser welding power.

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The quantitative measurements of partially melted zone for the fiber laser beam-welded superalloys are provided in Table 5. The results show that increasing the laser power and welding velocity does not affect the length of PMZ. So, the PMZ length is a function of only welding heat input. The quantitative results are plotted in Fig. 12. The curves indicate that the length of PMZ increases when the Nb content of the superalloys increases. In other words, the phases enriched in Nb display stronger susceptibility to liquation than do Ti-rich phases. This can be understood as the liquated phases exhibit lower temperature (T5) in the HAZ of the Alloy 5, and higher temperature (T1) in the HAZ of Alloy 1. Therefore, Alloy 1 and Alloy 5 display the highest and the lowest resistance to the hot liquation cracking, respectively.

3.4. Microhardness evaluation of the welds, heat-affected zones, and base metals Fig. 13 illustrates the Vickers microhardness for the welds, base metals and heat-affected zones of fiber laser beam-welded superalloys. The measurements demonstrate that changing the welding parameters (welding velocity and laser power) with constant heat input does not influence the hardness level of the weldments. Thus, the hardness is a function of only welding heat input. As seen, the welds of Alloy 1 and Alloy 5 show the maximum and minimum hardness values, respectively. The microhardness of the welds declines when the Nb content increases and Ti content decreases in the superalloy composition. The microhardness of the base metals follows the same order as that of the weld metals. Alloy 5 base metal shows very high hardness in comparison to its weld metal and heat affected zone due to formation of γ″-Ni3Nb strengthening precipitations. γ″-Ni3Nb precipitations have not been formed in the weld metal and these have decomposed in the HAZ of Alloy 5.

Fig. 6. Weld microstructure of fiber laser beam-welded newly-designed superalloys in different powers and welding velocities: (a) Alloy 1, P¼ 400 W, V ¼ 40 mm s  1, (b) Alloy 1, P ¼1000 W, V ¼ 100 mm s  1, (c) Alloy 2, P ¼400 W, V ¼40 mm s  1, (d) Alloy 2, P¼ 1000 W, V ¼100 mm s  1, (e) Alloy 3, P¼ 400 W, V ¼ 40 mm s  1, (f) Alloy 3, P¼ 1000 W, V¼ 100 mm s  1, (g) Alloy 4, P¼ 400 W, V ¼40 mm s  1, (h) Alloy 4, P¼ 1000 W, V¼ 100 mm s  1, (i) Alloy 5, P ¼400 W, V ¼40 mm s  1 and (j) Alloy 5, P ¼1000 W, V ¼ 100 mm s  1.

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Fig. 7. Heat-affected zone microstructure of fiber laser beam-welded newly-designed superalloy 1 in different powers and welding velocities: (a) P¼ 400 W, V ¼ 40 mm s  1, (b) P¼ 1000 W, V ¼ 100 mm s  1, (c) EDS analysis of the HAZ liquated regions: Cr- and Mo-rich eutectics, and (d) EDS analysis of the HAZ liquated regions: Ti-rich eutectics.

Fig. 8. Heat-affected zone microstructure of fiber laser beam-welded newly-designed superalloy 2 in different powers and welding velocities: (a) P¼400 W, V¼40 mm s  1, (b) P¼ 1000 W, V¼ 100 mm s  1, (c) EDS analysis of the HAZ liquated regions: Ti- and Nb-rich carbides, and (d) EDS analysis of the HAZ liquated regions: Nb- and Mo-rich eutectic structures.

Fig. 9. Heat-affected zone microstructure of fiber laser beam-welded newly-designed superalloy 3 in different powers and welding velocities: (a) P¼ 400 W, V¼40 mm s  1, (b) P¼1000 W, V¼ 100 mm s  1, (c) EDS analysis of the HAZ liquated regions: Ti- and Nb-rich carbides, and (d) EDS analysis of the HAZ liquated regions: Nb- and Mo-rich eutectic structures.

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Fig. 10. Heat-affected zone microstructure of fiber laser beam-welded newly-designed superalloy 4 in different powers and welding velocities: (a) P ¼400 W, V ¼ 40 mm s  1, (b) P¼ 1000 W, V ¼100 mm s  1, and (c) EDS analysis of the HAZ liquated regions: Nb- and Mo-rich Laves eutectics.

Fig. 11. Heat-affected zone microstructure of fiber laser beam-welded newly-designed superalloy 5 in different powers and welding velocities: (a) P¼ 400 W, V ¼ 40 mm s  1, (b) P¼ 1000 W, V ¼100 mm s  1, and (c) EDS analysis of the HAZ liquated regions: Nb-rich δ intermetallics.

4. Conclusions In the current research, the modern fiber laser beam welding of newly-designed precipitation-strengthened nickel-base superalloys using various welding parameters with constant heat input has been investigated. The results are summarized as follows:

 With increasing welding power and velocity at the constant  Fig. 12. Maximum length of partially melted zone and minimum temperature of partially melted zone in the fiber laser beam welded superalloys versus the superalloy designation.



heat input, the key-hole penetration mode is achieved and the penetration-to-width ratio of the welds is improved. With increasing welding power and velocity at the constant heat input, the porosity formation has been increased. This is not a function of superalloy composition. The solidification cracks are not seen in the welds, and the weld of Alloys 1 and 5 shows significant microsegregation.

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Fig. 13. Microhardness measurements of base metals, weld metals and heat-affected zone in the fiber laser beam welded superalloys versus the superalloy designation.

 None of the alloys show liquation cracking, but the liquation





phenomenon has taken place in all the HAZs. The PMZ length is a function of heat input, only. The length of PMZ increases when the Nb content of the superalloy composition increases. Changing the welding parameters does not alter the weldment microhardness. Thus, the hardness is a function of heat input only. The hardness of the welds increases when the Ti content of the superalloy increases. Based on the achieved results, 400 and 1000 W fiber laser powers with velocity of 40 and 100 mm s  1 are proposed for hot crack-free welding of thin sheet of precipitationstrengthened nickel-base superalloys.

Acknowledgment The authors thank the financial support of MAPNA Group under contract No. RD-THD-89-02. Professor Barbara Previtali, the director of SITEC—Laboratory for Laser Applications of Politecnico di Milano, and her team members: Mr. Daniele Colombo and Mr. Bruno Valsecchi are gratefully acknowledged for their laser facilities. References [1] Dupont JN, Lippold J, Kiser S. Welding metallurgy and weldability of nickelbase alloys. New Jersy: Wiley and Sons Inc. Publication; 2009. [2] Naffakh H, Shamanian M, Ashrafizadeh F. Dissimilar welding of AISI 310 austenitic stainless steel to nickel-based alloy Inconel 657. Journal of Materials Processing Technology 2009;209:3628–39. [3] Qian M, Lippold JC. Liquation phenomena in the simulated heat-affected zone of alloy 718 after multiple postweld heat treatment cycles. Welding Journal 2003:145–50. [4] Ojo OA, Richards NL, Chaturvedi MC. Contribution of constitutional liquation of gamma prime precipitate to weld HAZ cracking of cast Inconel 738 superalloy. Scripta Materialia 2004;50:641–6.

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