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V.T. Gaikwad, M.K. Mishra, V.D. Hiwarkar, and R.K.P. Singh, Microstructure and mechanical properties of friction welded carbon steel (EN24) and nickel-based superalloy (IN718), Int. J. Miner. Metall. Mater., 28(2021), No. 1, pp.111-119. https://dx.doi.org/10.1007/s12613-020-2008-1
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Continuous-drive rotary friction welding was performed to join cylindrical specimens of carbon steel (EN24) and nickel-based superalloy (IN718), and the microstructures of three distinct weld zones—the weld interface (WI)/thermo-mechanically affected zone (TMAZ), the heat-affected zone (HAZ), and the base metal—were examined. The joint was observed to be free of defects but featured uneven flash formation. Electron backscatter diffraction (EBSD) analysis showed substantial changes in high-angle grain boundaries, low-angle grain boundaries, and twin boundaries in the TMAZ and HAZ. Moreover, significant refinement in grain size (2–5 μm) was observed at the WI/TMAZ with reference to the base metal. The possible causes of these are discussed. The microhardness profile across the welded joint shows variation in hardness. The changes in hardness are ascribed to grain refinement, phase transformation, and the dissolution of strengthening precipitates. The tensile test results reveal that a joint efficiency of 100% can be achieved using this method.
Friction welding (FW) is an economical and environment-friendly solid-state joining process that is applied to join similar as well as dissimilar ferrous and nonferrous materials [1]. In this technique, under pressure, relative motion is introduced between the contact surfaces, leading to an increase in the interface temperature due to frictional heat. After the appropriate temperature is attained, rotational motion is stopped, and forging pressure is exerted followed by self-cooling [2]. The microstructure and the subsequent mechanical properties of friction welded joints are influenced by the applied pressure, friction and upset time, and rotational speed. Among recently introduced joining techniques, FW is attractive because of its lower cycle time, higher productivity, and better product quality [3].
To reduce cost, the properties of two different materials are combined for many engineering applications. Joints of carbon steel and nickel-based superalloy are used in the aerospace, automobile, shipping, and petrochemical industries [4−5]. The literature shows that FW has been effectively applied to join dissimilar materials, such as austenitic stainless steels to carbon steels [6−8], nickel-based superalloys with titanium alloys [9], and titanium alloys with stainless steel [10]. Moreover, some researchers have used this technique to join similar materials such as carbon steels, aluminum alloys, and nickel-based superalloys [1, 11−13]. Sahin et al. [13] joined AISI 1040 steel and reported grain size refinement at weld interface (WI), which could be favorable for the improvement of tensile strength and hardness of weld joint. Anitha et al. [14] observed that with increasing rotational speed and friction pressure, grain refinement occurred due to the enhancement of dynamic recrystallization (DRX). Khidhir and Baban [15] achieved 90% joint efficiency (discussed in Section 3.3.2) in joining AISI 1045 medium-carbon steel and AISI 316L austenitic stainless steel by FW methods. However, Kimur et al. [16] obtained 100% joint efficiency by joining a superalloy and heat-resistant steel by varying FW parameters.
Using the regular fusion welding methods to join Inconel alloys is challenging because of the occurrence of porosity, inclusion, cracking, and microstructural segregation. This problem can be reduced by the FW technique [4, 12]. Moreover, to further improve the weld joint strength of IN718 superalloy, post-weld heat treatment (PWHT) is suggested by Wang et al. [17]. Lalam et al. [5] observed that PWHT is beneficial to improve the joint strength but deteriorates the fracture toughness due to the precipitation of complex phases such as Ni3Al4, Fe3Mn7, and Al0.9Ni1.1. Damodaram et al. [18] joined IN718 rods under different heat treatment conditions and studied the effect of PWHT on the microstructure and mechanical properties of the weld joint. They concluded that after FW, double aging cycles (heating to 720°C for 8 h followed by furnace cooling up to 620°C and holding at 620°C for 8 h followed by air cooling) was the best method to restore the parent material properties.
Though FW is a solid-state joining process, some researchers have reported its negative impact on the joining of some dissimilar materials. During the FW of carbon steel and stainless steel, Ma et al. [6] detected carbides layer consisting of CrC and Cr23C6 at the WI, which are detrimental to mechanical properties. Similarly, Paventhan et al. [7] observed the formation of intermetallics such as Cr23C6 and Cr7C3 at the WI of medium-carbon steel and austenitic stainless steel, which deteriorate the fatigue strength. However, Madhusudhan Reddy and Venkata Ramana [19] tried to insert a nickel interlayer between high-strength maraging steel and low-carbon (EN24) steel to reduce the formation of intermetallic phases. Despite the development of new technology and the rapid automation in the industries, the FW of dissimilar materials and its implementation on the industrial level are still challenging tasks to researchers.
Considering the reports on the FW of dissimilar carbon steel (EN24) and nickel-based superalloy (IN718) are limited, the objective of the current research is to examine the microstructural evolution during FW and investigate the mechanical properties of the welded joint at room temperature.
EN24 steel and IN718 round bars were used as the base metal for FW. The diameter and length of each sample were 16 and 65 mm, respectively. The chemical compositions of both base metals are presented in Table 1. The EN24 steel rods were in annealed condition. Their microstructure consisted of spheroidite, i.e., sphere-like carbides in a ferrite matrix (Fig. 1(a)). As-received IN718 rods were solution-treated and aged. The microstructure of the IN718 consisted of strengthening precipitates and MC-type carbides in a γ (Ni) matrix (Fig. 1(b)). To eliminate the defect at the weld joint, the surface area of the joint was perfectly polished and then cleaned with acetone.
| Metal | Ni | Cr | Nb | Mo | Ti | Al | C | Mn | Si | Cu | Fe |
| EN24 | 1.53 | 1.29 | — | 0.25 | — | — | 0.41 | 0.61 | 0.26 | — | Bal. |
| IN718 | 53.60 | 17.82 | 5.05 | 3.01 | 0.66 | 0.45 | 0.08 | — | — | 0.04 | 16.91 |
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Friction welding was performed on a 120 KN upset force capacity continuous-drive FW machine (model FWT12), with all the FW process parameters precisely controlled by a computer numerical control system. The machine can operate with high precision and excellent repeatability, with spindle speed ranging from 800 to 2500 r/min. It can weld a maximum diameter of 32 mm. Welding was performed by holding the IN718 and EN24 samples in the rotating side and feeding side, respectively as shown in Fig. 2(a). Different trials were conducted by varying the friction pressure between 3000 and 7000 kPa, upset pressure between 6000 and 10000 kPa, and friction time between 2 and 10 s. The optimized parameters used for this study are presented in Table 2.
| Friction pressure / kPa | Friction time / s | Upset pressure / kPa | Upset time / s | Rotational speed / (r·min−1) |
| 5000 | 6 | 8000 | 3 | 1400 |
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After the FW, metallographic specimens were cut using a wire electrical discharge machine along the transverse direction of the weld joint. The cut samples were prepared by polishing and chemically etching according to the standard metallographic procedures. To reveal the microstructure of the IN718, Kalling’s reagent consisting of 10 g CuCl2, 200 mL HCl, and 200 mL ethanol was applied, whereas the EN24 side was etched using Nital. Microstructure analysis was performed using a scanning electron microscope equipped with a field emission gun. Chemical analysis was carried out using energy-dispersive spectroscopy (EDS). Samples for electron backscatter diffraction (EBSD) analysis were prepared by electro-polishing with an electrolyte of 200 mL perchloric acid, 100 mL 2-butoxy ethanol, and 700 mL ethanol. The parameters used for the electro-polishing were 24 V and 80 s. The EBSD scans were carried out at a step size of 0.4 μm using a Zeiss EVO18 scanning electron microscope. The grain boundary characteristics and grain size were determined using the Aztech HKL software. Tensile test specimens were prepared as per the ASTM E8 standard along the transverse direction to the weld joint, keeping the weld zone at the center of the gauge length. Tensile tests were carried out at ambient temperature, with a constant crosshead speed of 0.03 mm/s. Microhardness measurements were performed across the weld joint with a load of 500 g and dwell time of 10 s using a Zeiss microhardness testing machine.
The image of the welded joint is shown in Fig. 2(a). After careful visual inspection, the joint was confirmed to be without defects, indicating acceptable metallurgical properties. The reduction of axial length after FW is known as burn off, and the amount of material extruded is called flash [3]. As shown in Fig. 2(a), the total burn-off length was measured to be 9.0 mm; the reduction of lengths from WI to the EN24 side and IN718 side were 5.7 and 3.3 mm, respectively. The thermal conductivity of EN24 has been reported to be approximately 3.5 times higher than that of IN718 (IN718: 11.4 W·m−1·K−1 and EN24: 40.2 W·m−1·K−1) [4, 20]. Thus, more heat is accumulated at the friction interface of the IN718 side than the EN24 side because of the lower thermal conductivity, which leads to uneven temperature distribution across the WI. Furthermore, at a high temperature, IN718 has a higher tensile strength compared with steel. Therefore, the EN24 experienced more deformation and extruded more material in terms of flash at the WI. The inverse pole figure (IPF) of the welded joint is presented in Fig. 2(b).
The weldment can be broadly divided into three zones: (1) the WI and/or thermo-mechanically affected zone (TMAZ), (2) the heat-affected zone (HAZ), and (3) the unaffected base metal at both sides. The WI thickness is very narrow, around 30 μm. The microstructure in this region consists of fine grains with an intermixture of EN24 and IN718 materials. The microstructural variations of the different zones on the EN24 alloy and IN718 are discussed in detail in Sections 3.1.1 and 3.1.2, respectively.
The microstructures in different zones of the rotary friction welded carbon steel (EN24) sides are presented in Fig. 3. The microstructure of the base metal consisted of spheroidite carbide particles distributed uniformly in the ferrite matrix as shown in Fig. 1(a). In the HAZ, deformed grains of fine pearlite and ferrite were observed (Fig. 3(a)). However, in the WI/TMAZ, a martensitic phase was observed (Fig. 3(b)). Due to friction, the temperature of the WI increased and reached the austenitic temperature range. When the rotational motion followed by upset pressure is stopped, self-quenching occurs in the bulk materials, resulting in the phase transformation of austenite to martensite [5, 21].
The IN718 superalloy has good corrosion resistance and mechanical properties at elevated temperatures. The main strengthening precipitates for this alloy are γ′ [Ni3(Al,Ti)] and γ″ (Ni3Nb) [4]. Fig. 4 shows the microstructures of the IN718 side at various zones. The base metal (Figs. 4(a) and 4(a′)) comprised the γ matrix with fine γ′ and γ″ precipitates and a needle-shaped delta phase (δ) along grain boundaries. Moreover, titanium carbide and niobium carbide particles were dispersed in the matrix. The presence of these carbides and δ phase were confirmed by EDS point scanning (Figs. 4(a1), 4(a2), and 4(a3)), and the weight percentages of elements are presented in Table 3.
| Point | C | Ti | Cr | Fe | Ni | Nb | N | O | Mg | Al | Total |
| 1 | 13.21 | 1.51 | 6.28 | 7.17 | 52.13 | 19.69 | — | — | — | — | 100 |
| 2 | 7.39 | 45.41 | 0.56 | — | 0.83 | 8.27 | 20.01 | 15.64 | 0.79 | 1.1 | 100 |
| 3 | 30.39 | 7.22 | 0.73 | — | 1.05 | 57.48 | — | 3.13 | — | — | 100 |
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Next to the base metal, elongated grains were present in the HAZ (Figs. 4(b) and 4(b′)). However, the WI/TMAZ showed a significant reduction in grain size due to the occurrence of DRX during welding, as indicated in Figs. 4(c) and 4(c′). A temperature beyond 850°C will result in the dissolution of γ′ and γ″ hard precipitates and δ phases in alloy 718 [12]. As observed from the SEM image, these precipitates were dissolved at the WI/TMAZ, depending on the temperature. To study the elemental distribution across the WI, an EDS line scan was performed, and the results are presented in Fig. 5. A significant change in element distribution, mainly for Fe, Ni, Nb, Cr, and Ti, was observed.
Furthermore, the lines corresponding to these elements present large scattering near WI, which is a mechanically mixed zone of two dissimilar materials (EN24 and IN718).
The IPF maps and misorientation distributions for the EN24 side are shown in Fig. 6. From Fig. 6(a), it can be understood that the base metal of the EN24 consisted of comparatively coarse grains. The fractions of HAGB and LAGB in the base metal were observed to be 0.04 and 0.96, respectively, and they drastically changed in the HAZ to 0.63 and 0.37, respectively. The increase in the HAGB and decrease in the LAGB are associated with DRX. It has been well established that DRX occurs during FW [12, 22]. The critical strain for the recrystallization is provided by the plastic deformation produced by the upset force and the required temperature due to friction at the joining interface. However, the HAGB and LAGB fractions in the WI/TMAZ were found to be 0.36 and 0.64, respectively. The increase in the LAGB fraction compared with that of the HAZ may be due to the development of a martensite structure [23].
The IPF maps and misorientation distribution maps for the IN718 are given in Fig. 7. The base metal of the IN718 consisted of equiaxed grains with an average grain diameter of 20–22 μm. However, the HAZ showed slightly deformed grains whose size is close to that of the base-metal grains. Significant refinement in grain size up to 2–5 μm due to DRX was observed at the WI/TMAZ. From the misorientation distribution map, the HAGB and LAGB fractions were 0.94 and 0.06 in the IN718 base metal, respectively. However, in the HAZ, they were 0.78 and 0.22, respectively. Moreover, the WI showed a drastic change in grain boundary characteristics. The HAGB and LAGB fractions were observed to be 0.35 and 0.65, respectively.
The lower HAGB fraction and higher LAGB fraction in the HAZ and TMAZ compared with those in the base metal indicates that full static recrystallization did not take place in these zones. However, the percentage of twin boundaries (misorientation angle close to 60°) shows a reducing trend from the base metal to the WI/TMAZ. The fractions of the twin boundaries in the base metal, HAZ, and WI/TMAZ are estimated to be around 0.18, 0.16, and 0.02, respectively. The large reduction in twin boundaries at the WI/TMAZ occurred because this region underwent severe plastic deformation at a high temperature. It has been reported that the ideal twin–matrix relationship is destroyed during the deformation, which decreases the fraction of the twin boundaries [24−25].
The microhardness variation across the WI is illustrated in Fig. 8. The variation of microhardness across the joint is related to the microstructure formed in the various welding zones. In the WI/TMAZ zone of the EN24 side, the peak microhardness value was found to be HV 635, which is significantly higher than that of the base metals. The increase in hardness is due to the combined effect of a martensitic structure formation and refinement in grain size (2–3 μm).
Grain refinement has been reported to occur due to DRX when the material experiences plastic deformation at high strain rates under elevated temperatures [26−27]. The hardness in the HAZ exhibited a decreasing trend from the WI toward the base metal, which is attributed to the formation of ferrite and pearlite in the microstructure.
Furthermore, the precipitates of the nickel-based superalloy (IN718) strengthened the material; the hardness of the material depends on the fractions of the γ′ and γ″ precipitates in the microstructure [4]. The WI/TMAZ and HAZ had a lower hardness compared with the base metal. This could be due to the dissolution of strengthening precipitates γ′ and γ″. The base metal, which was not affected by heat and deformation, retained the hardness of the as-received parent metal, i.e., HV 440.
The yield strength (YS), ultimate tensile strength (UTS), elongation, and joint efficiency of the welded joint are presented in Table 4. The weld samples showed YS that were almost equivalent to that of the EN24 base metal. However, the weld samples exhibited a lower ductility, which is due to the grain size refinement and martensite formation in the weld zone. Khidhir and Baban, [15] also reported lower ductility of the welded joint of dissimilar materials of AISI 1045 carbon steel and 316L austenitic stainless steel. The efficiency of the welded joint can be calculated as follows [15]:
| Sample | Yield strength (YS) / MPa | Ultimate tensile strength (UTS) / MPa | Elongation / % | Joint efficiency (η) / % |
| Base-EN24 | 675 ± 7.0 | 767 ± 15.0 | 22 ± 1.5 | |
| Base-IN718 | 1188 ± 8.0 | 1374 ± 9.0 | 24 ± 3.0 | |
| Weld sample | 681 ± 5.0 | 772 ± 10.0 | 14.8 ± 1.0 | ~100.0 |
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| Jointefficiency(η)=TensilestrengthofweldjointTensilestrengthofsoftmaterials(EN24steel)×100% |
The UTS of the welded joint was around (772 ± 10.0) MPa, which shows a joint efficiency of ~100%. Therefore, the selected welding parameters are appropriate to generate friction heat and create plastic deformation at joints.
During tensile testing, failure occurred in weak or lower-hardness zones (base EN24). The SEM image of the fracture surface of the tensile-tested sample is shown in Fig. 9. The fractographs show the presence of dimples, indicating ductile failure.
The following conclusions can be drawn from the present work on the FW of dissimilar EN24 alloy steel and IN718 superalloy.
(1) The welded joint was observed to be free of defects, with the overall burn-off length being 9.0 mm. However, the flash width was uneven due to the difference in thermal conductivity and yield strength between the two metals.
(2) The EBSD results showed a significant refinement of grain size (2–5 μm) in the IN718 TMAZ. The grain boundary characteristics of the WI/TMAZ and HAZ were different from that of the base metal due to the combined effect of heat and severe plastic deformation. Moreover, the fraction of twin boundaries in the WI was lower than that of the base metal, and this can be ascribed to the severe thermo-mechanical deformation.
(3) The hardness variation across the weld joint is attributed to grain refinement, phase transformation, and the presence of strengthening precipitates. The peak hardness (HV 635) was observed toward the TMAZ of the EN24 side, due to the formation of a martensite structure and grain refinement, while the TMAZ of the IN718 exhibited a lower hardness (HV 385) compared with that of the base metal due to dissolution of strengthening precipitates. Moreover, the UTS of the welded joint was equivalent to that of the EN24 base metal.
The authors acknowledge Kalyani Centre for Technology & Innovation (KCTI), Bharat Forge Ltd, Pune, and Department of Scientific and Industrial Research (DSIR), goverment of India for providing financial assistance, library, and laboratory facilities. The authors are gratefulness to Review Committee of Bharat Forge Ltd, Pune to allow for publication. Also, authors are thankful to Vinayak Pawar and Santosh Dhage for their assistance in various aspects of the project. Finally, authors thank K. Chethan for valuable suggestion, guidance, and discussion.
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| [9] | Ji-xiong Liu, Ai-min Zhao, Hai-tao Jiang, Di Tang, Xiao-ge Duan, Heng-yong Shui. Effect of interface morphology on the mechanical properties of titanium clad steel plates [J]. International Journal of Minerals, Metallurgy and Materials, 2012, 19(5): 404-408. DOI: 10.1007/s12613-012-0571-9 |
| [10] | Guo-wei Chang, Shu-ying Chen, Qing-chun Li, Xu-dong Yue, Guang-can Jin. Lateral growth rate of δ phase during solidification of Fe-0.15wt%C-0.8wt%Mn steel [J]. International Journal of Minerals, Metallurgy and Materials, 2012, 19(4): 312-316. DOI: 10.1007/s12613-012-0556-8 |
Figures(9) / Tables(4)
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| Metal | Ni | Cr | Nb | Mo | Ti | Al | C | Mn | Si | Cu | Fe |
| EN24 | 1.53 | 1.29 | — | 0.25 | — | — | 0.41 | 0.61 | 0.26 | — | Bal. |
| IN718 | 53.60 | 17.82 | 5.05 | 3.01 | 0.66 | 0.45 | 0.08 | — | — | 0.04 | 16.91 |
DownLoad:
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| Friction pressure / kPa | Friction time / s | Upset pressure / kPa | Upset time / s | Rotational speed / (r·min−1) |
| 5000 | 6 | 8000 | 3 | 1400 |
DownLoad:
CSV
| Point | C | Ti | Cr | Fe | Ni | Nb | N | O | Mg | Al | Total |
| 1 | 13.21 | 1.51 | 6.28 | 7.17 | 52.13 | 19.69 | — | — | — | — | 100 |
| 2 | 7.39 | 45.41 | 0.56 | — | 0.83 | 8.27 | 20.01 | 15.64 | 0.79 | 1.1 | 100 |
| 3 | 30.39 | 7.22 | 0.73 | — | 1.05 | 57.48 | — | 3.13 | — | — | 100 |
DownLoad:
CSV
| Sample | Yield strength (YS) / MPa | Ultimate tensile strength (UTS) / MPa | Elongation / % | Joint efficiency (η) / % |
| Base-EN24 | 675 ± 7.0 | 767 ± 15.0 | 22 ± 1.5 | |
| Base-IN718 | 1188 ± 8.0 | 1374 ± 9.0 | 24 ± 3.0 | |
| Weld sample | 681 ± 5.0 | 772 ± 10.0 | 14.8 ± 1.0 | ~100.0 |
DownLoad:
CSV
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