A Review on Friction Stir Welding of Steels

Friction Stir Welding (FSW), as a solid-state welding process, can achieve the joining of metals below the melting point or in the plastic stage of the metals with the aid of a non-consumable tool [1]. Compared with the welds made by fusion welding processes, the friction stir welds have a more homogeneous grain structure and better mechanical properties like tensile strength, hardness, and toughness [2]. Friction stir welded specimens exhibit better corrosion resistance properties due to the homogeneous microstructure and narrow heat-affected zone. The feature of joining metals in the semisolids state helps this FSW method to join metals with different melting points like aluminium and copper. Nowadays, FSW has become one of the prevailing joining processes of aluminium alloys in the industry [3‐5]. While considering hard metals like steel, FSW still does not accomplish the same feasibility in aluminium alloys [6, 7]. The limitations of friction stir welding on steels are as follows: (a) a very high durable tool is required for welding steels, (b) the temperature produced by the tool pin and shoulder will not be sufficient to plasticize the metals, (c) welding speed cannot be attained as good as on aluminium alloys due to the high hardness of steels, (d) tool damage rate is very high, and (e) high flow stress is maintained by the hot steels while conducting FSW process, and which causes high contact stress and severe tool degradations [8]. Tool material like commercial pure Tungsten (Cp- W) or Polycrystalline Cubic Boron Nitride (PCBN) tools can survive these problems. Still, it increases the welding cost by eight times or even more while comparing to fusion welding.

Although the FSW tools for steel are expensive, friction stir welding of steel has numerous advantages compared to conventional fusion welding methods. Some of them are as follows: (a) the heat-affected zone created by fusion welding is large, and this eventually decreases the strength and quality of the weld, (b) fusion welding requires a filler metal to join the base metals, where FSW does not require any filler metals at all, (c) to eliminate hydrogen embrittlement or porosity like defects, fusion welding requires shielding gas which will pollute the atmosphere highly, while FSW does not require any kind of shielding gases, (d) fusion welding arcs emit intense visible, ultraviolet, and infrared radiation which is very hazardous to the welder and nature, whereas FSW does not produce any harmful rays or FSW is a 100% eco-friendly joining process [9‐11]. Moreover, fusion welding of martensitic steels is complicated, but for FSW, it is pretty easy.

To improve the welding efficiency and quality of steel, hybrid friction stir welding was invented. An additional heating source will combine with the conventional FSW for in-situ heating in hybrid friction stir welding. The best examples for such welding are laser-assisted FSW (LA-FSW) and induction assisted FSW (IA-FSW) [12, 13]. These methods help to in-situ heat the metals, and by this way, a higher welding speed will achieve; the softening of metal before plunging the tool helps reduce tool damage rates, and tools like high-speed steel tungsten carbide were capable of welding steels [14]. In these methods, the temperature requires to weld steel ranges from 500 °C to 1100 °C. The microstructure analysis of such joints proved that the specimen has a fine grain structure and better mechanical properties than fusion welding [15]. These researches show that FSW on steel became much cheaper and commercial, like FSW on aluminium alloys.

This article reviews the current research and advancements in friction stir welding of steels and determines all the constraints faced in this present scenario.

The tool materials widely used for FSW are high-speed steel (HSS). Soft metals like aluminium, copper and their alloys can often be welded by this kind of tool [16‐18]. For overcoming tool wear, process parameters optimization is mandatory [19]. In welding steels, HSS tools are incapable because they cannot produce the required temperature, withstand the elevated temperature of about 950 °C, and resist the abrasive wear [20‐22]. The tools for joining steels must tolerate all kinds of wear, have high strength and hardness, and have fracture toughness. Considering all these properties, the tool materials used nowadays are Commercially pure Tungsten (Cp- W) and Poly Crystalline Cubic Boron Nitride (PCBN).

The vital issue of fusion welding is its limitation to weld dissimilar metals and the heavy heat affected zone created in the joining process. Compared to fusion welding, the heat generation is significantly less in FSW. FSW can overcome the problems encountered by fusion welding due to its solid-state welding feature.

11 pct Cr ferritic/martensitic steel (PNC-FMS) and 316 austenitic stainless steel were successfully welded using the FSW process. AISI 316 was kept at the advancing side and the PNC-FMS at the retreating side for this welding. A defect-free joint was produced, but the two steel plates were not mixed well but interleaved a sharp zigzagging interface in the stir zone [143]. Besides, the robust interface did not act as a fracture site when a micro tensile test was conducted for the stir zone alone. The mechanical property revealed by the specimen was much better than the base metals; this indicates that FSW can be feasible for sound welding in dissimilar metals, too [144, 145]. Figure 13 shows the transverse cross-section of the joints taken by optical microscopy, revealing that this joint does not have any weld defects, like tunnel defects, cracks and voids.

The scanning electron microscope image and the elemental mapping done by energy-dispersive x-ray spectroscopy (EDS) images are given in Figure 14. The lighter contrast in the EDS mapping images means the higher content, and the darker means, the lower content, respectively [146]. The specimen selected for elemental mapping was prepared at a tool rotation of 100 r/min. The Fe in the 316 stainless steel side is lower when the Cr and Ni content is lower on the PMC-FMS side [147‐151]. These distributed elements are separated along with the weld interface.

Figure 15 shows the magnified SEM image of the stir zone. There are no voids or reaction layers visible in it, and the interface’s sharpness is very well observed [152]. A gradual change in Cr and Ni concentrations has occurred over the interface due to the diffusion of alloying elements during friction stir welding. In this experiment, the tool rotation speed of 100 r/min and 150 r/min exhibited similar results. The only difference was in the width of diffusion layers, and it was more prominent at the tool rotation speed of 150 r/min [153]. The results have proven that intermetallic bonding was successful in PNC-FMC and 316 stainless steel joints.

During friction stir welding of F82H to 304 stainless steel, the shift in tool plunging prevented the intermixing in the nugget zone [154, 155], so the exact extent of mixing was in the stir zone not precise. Moreover, the welding conditions like tool profile and welding speed predominantly influence the nugget zone’s grain formation [156‐159]. Therefore, this research shows that any change in the plunging area may lead to kissing bonding or root defects.

In dissimilar welding of steel with other metals like superalloys [160], the melting point of these alloys is somewhat close. The super duplex steel SAF 2507 was friction stir welded with a nickel-based superalloy Incoloy 825. A defect-free weld was fabricated by placing SAF 2507 on the advancing side and Incoloy 825 on the retreating side [161, 162]. The microstructure analysis has shown that the thermo-mechanically affected side was from deformed austenite−ferrite grains of the SAF 2507 region and the elongated austenite at the Incoloy 825 region. Due to the subsequent refinement of the grain structure and dynamic recrystallization, the nugget zone shown the highest hardness.

In the medium carbon steel welding, carbon and chromium presence significantly affect the weld region’s microstructure formation. The chromium content enhancement in the stir zone transformed the pearlite structure into a martensite structure [163, 164]. This transformation would increase the tensile strength and ductility of the joint. The carbon and chromium content also reduces the carbide precipitation in the stir zone due to martensite's auto tempering [165].

Due to the supercilious comprehensive properties of the reduced activation, ferritic/martensitic (RAFM) steels are broadly used in nuclear reactors as structural materials [166, 167]. An investigation was conducted to join the reduced activation ferritic/martensitic steel by friction stir welding. The quantitative relationship with microstructure generated in the stir zone and tensile strength was examined [168‐171]. Moreover, the effects of post-weld tempering treatments were also studied. For this post weld tempering treatment, three different temperatures were adopted, i.e., 720 °C, 760 °C, and 800 °C, respectively [172]. The microstructure evaluation was conducted by using transmission electron microscopy (TEM), Electron backscatter diffraction (EBSD) and the Thermo-Calc Calphad software [173]. The matrix’s precipitate number density, martensite lath width, geometrically necessary dislocation (GND) density, and solid equilibrium solubility of alloying elements were quantitatively analyzed.

The transmission electron microscope image of the RAFM steel is given in Figure 16(a) and 16(b). In Figure 16(a), a distinctive tempered martensitic microstructure is visible in it, and coarse M23C6 particles are seen in the martensite lath boundaries. In Figure 16(b) the inner area of the martensite laths MX carbonitrides are distributed uniformly [174‐176]. Based on the interpretation, the size of the M23C6 particles ranges from 90‒230 nm, and the size of the MX carbonitrides varies from 15‒60 nm. In the intra-lath region, a high density of free dislocation is visible. The energy dispersive spectroscopy graphs of two different areas with different sizes marked as A and B in Figure 16(b) are given in Figure 16(c) and (d), respectively [177]. Cr and W rich M23C6 carbide are visible in the coarse particle at the A region because the A region mainly has Cr, W and C elements. Besides, the B region shows its V, Ta and C, N-type carbonitrides due to the chemical composition in this small region are V, Ta. C and N.

The stir zone TEM microstructure of the specimen after welding and after tempering is given in Figure 17. Due to the martensitic transformations in the FSW process, the martensite lath is visible in the stir zone (Figure 17(a)). The results show that while increasing the tempering temperature, the average width of martensitic lath increases exponentially [178]. Figure 17(b) shows that the specimen treated with tempering temperature 720 °C and the non-tempered specimen has appropriately the same average width of martensitic lath. The dislocation migration due to high-temperature recovery caused dislocation tangles and networks at a tempering temperature of 720 °C. Moreover, fragmented subgrains were formed inside the martensite laths when the tempering temperature was increased to 760 °C (Figure 17(c)). At tempering at 800 °C, it was nearly impossible to identify martensite laths, while an equiaxed grain or subgrains were found (Figure 17(d)).

Figure 18 shows the martensite laths width or the subgrains, determined using the line intercept method. The average width of martensite lath width was 441 nm when the stir zone was tempered at 760 °C. If the stir zone was tempered at 800 °C, the average width of martensite lath was 602 nm, which exceeds that of the base metal (493 nm) [179].

Figure 19 shows the high magnified transmission electron microscope images of the stir zone, which helps identify the precipitate distribution of the stir zone. Figure 19(a) reveals the dissolved MX carbonitrides dispensed within the martensite lath [180, 181]. The peak temperature attained by the stir zone ranged from 973 °C to 1139 °C. Due to this temperature, M23C6 carbides were hard to find. Thermo-Calc software with the thermodynamic database TCFE9 was used to analyze the equilibrium phases at various temperatures, which helps investigate the evolution of participants. The derived results are given in Figure 20.

In microstructure analysis, the phases like Laves phase and the Z phase are not visible in the base metal or stir zone of the RAFM steel, so they were omitted from the calculation. Figure 20 shows that below 886 °C, M23C6 carbides were formed. When the temperature was above 988 °C, the MX carbonitrides were dissolved entirely [182], indicating that MX carbonitride withstands higher temperatures than M23C6 carbides. Figure 20 shows the Thermo-Calc calculated data as well. Figure 19(b), (c) and (d) show that the thawed M23C6 and MX phases were reappeared from the supersaturated solid solution of the martensite matrix after PWTT [183, 184]. In contrast, the friction stir welding of super duplex steel SAF 2507 with a nickel-based superalloy Incoloy 825 indicates that vital microstructure deformation is happened due to the dynamic recrystallization on the dissimilar metal interface, and this will affect the yield strength of the joint.

Multi-material design or multi-material joining have a colossal requirement since joining dissimilar metals with different physical properties and chemical properties can provide a synergic effect like corrosion resistance, high weight reduction, and high impact & yield strength. These are the results gained from the hybrid uses of such dissimilar metals [185]. With the aid of friction stir welding, steel and aluminium were joined, and this solid-state welding avoids the formation of thick intermetallic compound layers in steel/aluminium joints. The optimized process parameters and a cylindrical tool geometry were the keys for this dissimilar joining process [186]. Three different tools were used for this experiment. For tool-1, shoulder diameter was 25 mm, pin length was 2.85 mm, and pin profile was tapered cylindrical with a 6‒8 mm diameter. For tool-2, shoulder diameter 18 mm, pin length 2.85 mm, tapered cylindrical pin with diameter 5‒7 mm. For tool-3, shoulder diameter 13 mm, pin length 2.85 mm, and cylindrical pin profile with diameter 5 mm [187, 188].

The bead geometry and microstructure analysis were carried out using an optical microscope and field emission scanning electron microscope (FESEM) for analyzing the intermetallic compounds (IMC) layer. The phase and composition of the IMC layer were investigated by the XRD method [189]. Figure 21 shows the field emission scanning electron microscope images of the joint interface at higher magnifications.

Figure 21(a) presents the magnified view of the joint interface made by tool-1. It reveals a serrated type of material, which formulae an IMC layer with fragmented particles. In Figure 21(b), the interface of weldment produced by tool-2 also shows such a damaged IMC layer. A tinny and uninterrupted IMC layer was visible in the weld interface fabricated by tool-3, as shown in Figure 21(c). While exposed to heat, the metallurgically mismatched steel/aluminium joint formed a brittle and hard IMC layer on the close contact region [190]. The joint strength is decreased according to the increase in the composition and thickness of this IMC layer.

The FESEM images also reveal that a unique IMC layer of somewhat different contrast from the parent metal is observable. Also, Figure 21 shows that the IMC layer’s thickness varies from 1 μm to 4 μm. This IMC layer variation occurs due to the different diffusion temperatures formed in other regions and the different kinetics of diffusion of materials [191, 192]. For all three joints, the thickness of all the weld interfaces was measured and shown in Figure 21(a1), (b1) and (c1). 10 points were taken for average thickness. The IMC layer thickness gained by tool-1 is comparatively higher than those formed by the other two tools. The continuous detachment of debris always limits the IMC layer growth. Consequently, for tool-1 and tool-2, the IMC layer area and the detached fragments nearby get amplified in the stir zone [193, 194]. As evident from Figure 21, the thickness of the IMC layer is about 2.5 μm.

With the aid of energy-dispersive X-ray spectroscopy (EDS), the chemical line analysis was conducted along the thickest region of the intermetallic layer, as shown in Figure 22(a) and 22(b). This EDS image clearly shows the coexistence of Al and Fe in the intermetallic layer in the line analysis curve. To check the existence of various elements, a chemical analysis was initiated for a selected point (Figure 22(c)) [195]. The element analysis method analysed the atomic percentage and weight of different aspects, given in Figure 22(d). The result revealed that the atomic percentage of the aluminium atom was around 73.42%, and the atomic ratio of iron was 25.45%. It is challenging to analyze the intermetallic phases of various composites in this complex system.

Along with the detached steel fragment, elemental analysis was conducted, shown in Figure 22(e) [196]. Reinforced particles were formed in the aluminium matrix due to the congestion of steel fragments in the nugget zone. In Figure 22(f), an IMC layer is visible around the steel debris with a thickness of 5 μm, thicker than the actual interface of aluminium/ steel [197, 198]. Initially, the detached fragments were part of interfacial steel and an IMC layer formed towards the aluminium side of the exposed surface. While examining the nugget zone, the debris gained a suitable diffusion time due to the higher heat generation, leading to thickening of the IMC layer.

XRD analysis was conducted for all the fracture tested specimens joined by three tools. The XRD analysis is given in Figure 23.

The XRD analysis showed that the joint fabricated by the tool-1 had the presence of Fe2Al5 beside aluminium and iron. For the tool-2 welded specimen, a more or less similar result was gained. The presence of aluminium metal was visible in the fracture sample made by tool-3. It was evident that the failure in this specimen was due to the aluminium metal and not the IMC layer [199].

With the aid of FESEM, with varying magnifications, the fractured tensile samples were examined, as shown in Figure 24. Figure 24(a) was the joint fabricated by tool-1. For improved clarity, a magnified view was given inside the box [200]. The image portrays that the failure has happened in the down portion of the IMC layer, and a small number of indiscretions were visible in the upper part of the fracture, and the remaining region was almost flat. Due to the high heat generated by the tool plunge depth, an extreme softened aluminium was formed in the upper area of the specimen (Figure 24(a)), which led to the fracture. Some unevenness was visible in the upper area of the fractured specimen prepared by the tool-2, but the bottom area was primarily flat. Figure 24(b) shows the occurrence of dimples in the magnified view, and these dimples caused ductile failure [201, 202]. Due to the tool spindle speed, the aluminium particles moved upwards along with the tool shoulder over the steel region, leading to the failure of the upper part of the fracture surface, and brittle fracture happened in the bottom.

Nevertheless, the steel and aluminium created an almost flat interface. Figure 24(c) shows the fracture surface of the weldment prepared by tool-3. A ductile failure was visible in the image, and in the magnified view, a web-like arrangement and dimples were visual, which reveals that the failure was happened by the ductile way. Also, some irregular surfaces were visible at the bottom of the specimen, which refers to the interaction of two metals [203, 204]. The aluminium presence in the fracture surface was visible in the XRD graph of the fracture specimen prepared by tool-3, and steel defragmentation was not due to the proper bonding. The results reveal that a tool causing maximum material movement can provide a better bonding surface and join steel with soft metals like aluminium. The steel grains should be minimized to obtain sound joint strength [205, 206].

Dual-phase (DP) steel 700 was welded successfully by using friction stir welding. The texture development and microstructure formation were examined thoroughly. FESEM analyzed the stir zone microstructure of the DP 700 weldment, as given in Figure 25.

The optical micrograph and the FESEM micrograph of the DP steel stir zone are shown in Figure 25(a) and (b). The image portraits that the stir zone combination of fine martensite, ferrite and retained austenite microstructure. Figure 25(c) and (d) show the stir zone's IPF map and phase map [207]. The grain size in the stir zone had an average length of 1.3 ± 0.5 μm, which is comparatively smaller than that in the base metal. Moreover, the 3.8% of retained austenite dispersed in the ferrite matrix and visible in the phase distribution map (Figure 25(d)). Besides, the austenite's prolonged crystal defect enhanced its constancy against martensitic evolution, resulting in an improved level of retained austenite in the microstructure [208, 209]. Nevertheless, the martensite and ferrite have a BCC structure, so it is hard to identify duplex steels. However, in the IFP map, this phase is seen as a cluster of comparatively arbitrarily oriented points.

Furthermore, the Kernal average Misorientation (KAM) maps were also interpreted to examine the validity and accuracy of the threshold. Significant differences are detected in the intermediate level of distortions in different grains, a grain distinct related to the minor boundaries of 50, and the most petite pixels of 4 is significant [210].

Based on the above Kernal average Misorientation map approach, the separation of martensite and ferrite is given in figure 26. The volume fraction was estimated to be 11.3% and 84.9% for martensite and ferrite, respectively. This volume fraction was calculated from many SEM images analyses [211]. The martensite and austenite elements were eliminated, and the IPF map of the ferrite phase was plotted in Figure 26(d). Figure 26(e)‒(f) show the grain boundaries map and misorientation angle distribution. It is visible that primarily all the HABs established closed structures, and equiaxed ferrite grains were also visible [212‐214]. The yellow arrow in Figure 26(e) represents the structure free of LABs, which symbolize that recrystallization had been finished. During the dynamic recovery process, steady accumulation and reordering of grains dislocations led to the formation of grains having LABs. The green arrows and red arrows in Figure 26(e) represent the adequate extent of dynamic recovery and inadequate dynamic recovery during the process, respectively [215]. Figure 26(f) illustrates that most of the LABs vary from 20 to 50. The research found that the recovery mechanism for the ferrite phase in the stir zone was continuous dynamic recrystallization and discontinuous dynamic recrystallization. Due to the lower strain in the thermomechanically affected zone, geometric dynamic crystallization was also involved in the ferritic phase recovery [216]. The primary recovery mechanism for ferrite in the thermomechanically affected zone and stir zone were continuous dynamic recrystallization.

Friction stir welding on 12Cr oxide dispersion strengthened (ODS) steel was successfully executed. The results have shown that the grain size in the stir zone was smaller than the base metal. The dynamic recrystallization in the stir zone and the thermomechanically affected zone was analyzed by electron backscattering diffraction (EBSD) results. It was observed that a changeover from low angle boundaries (LAB) to high angle boundaries (HAB) was convoyed by the dynamic recrystallization process [217].

The 12Cr-OSD steel weldments IPF maps were taken from different areas of the joint, as shown in Figure 27. The different colours and patterns in Figure 27(f) represent different crystallographic orientations.

The average grain size of the various zones was measured by the intercept method, and Figure 28 portraits the histogram of grain size distribution. Figure 28(a) shows the number fraction of grain size in the base metal. The grain size in base metal was 1249 nm, where the grain size in the heat-affected zone was 1121 nm, which was comparatively smaller than the base metal (Figure 28(b)); this is due to the recrystallization that happened in the metal due to heat generated by FSW process [218‐220]. The grain structure found in the stir zone was vastly different from the base metal and HAZ. The grain size in the stir zone was 358 nm, and the refined grains were uniformly distributed due to the dynamic recrystallization in the stir zone, as shown in Figure 28(c). The breakage of grains in the stir zone resulted in a refined grain structure. The heat-affected zone and thermomechanical affected zone grains did not break much and resulted in 878 nm and 1245 nm, respectively [221, 222].

Figure 29 shows the 12Cr-ODS steel joint’s EBSD grain boundary misorientation angle distribution in base metal, stir zone and thermomechanical affected zone. The blue line showed the misorientation angle of more than 150, i.e., the high angle boundary. The misorientation angle less than 150, i.e., the low angle boundary (50 to 150), was portraited in green, and the red line showed those with misorientation angle20 to 50. Figure 29(a) shows that the fraction of LABs was around 49.9 % and the elongated grains enclosed by the vast number of HABs. HAB formation increased to 86%, the metal flow and mechanical force highly promoted the recrystallization in the stir zone [223, 224]. Figure 29(b) shows that the grain boundaries of the stir zone were more refined, and the grain was distributed evenly compared to the base metal. Equiaxed grains and the LAB distributed besides the flow track were noticeable in the TMAZ region due to the recrystallization process, as shown in Figure 29(c). The researchers concluded that 49.4%, 85.6% and 49.6% were the recrystallization percentages in the heat-affected zone, stir zone and thermomechanical affected zone, respectively [225].

Implementing FSW for high melting point metals like steel is quite challenging due to the tool wear. The replacement of tools in a short span of welding makes the process costly. One of the best methods to attain a prolonged tool life is to adopt preheating or in-situ heating techniques to soften the weld region before the tool enters. This softening of the weld region will reduce the wear rate of the tool, and therefore the tool life will increase. This method is called Hybrid Friction Stir Welding.

Hybrid FSW means an additional heating source is attached with the conventional FSW process for preheating or in-situ heating the metal plates before the tool plunging [226‐228]. This method helps to plasticize the tool plunging area of the metals, and through this, the plunging and tool movement will be much easier. Figure 30 shows the schematic diagram for a hybrid FSW process.

This method helps to reduce the tool damages due to high axial force while plunging. Through this method, FSW of steels can take over by using HSS or tungsten carbide tools. The welding cost will be much cheaper through this technique than the PCBN tool used in the FSW process [229]. The widely used preheating or in-situ heating techniques are laser heating, ultrasonic vibration and induction heating. Laser heating can be applicable for magnetic and non-magnetic alloys, while induction is more capable in ferromagnetic steels [230‐232].

Friction stir welding of steels reached a certain height recently. Many types of research have been conducted to find better combinations like tool material, tool profile, welding speed, tool rotation speed, axial force, and tool tilt angles. This research shows that better steel joints can be fabricated by using friction stir welding [248]. While considering the structure changes, the primary factor affecting heat generation is welding and cooling rate. A fine-grain structure can be achieved in the weld region by controlling the heat generation and cooling rate after welding [249]. This method will enhance the tensile as well as the toughness of the joints. While comparing with the conventional fusion welding techniques, this friction stir welding has many advantages. The pollution that occurs while friction stir welding is nearly nil; porosity and hydrogen embrittlement should not happen while using FSW. Moreover, the heat-affected zone is comparatively less, and the mechanical and corrosive properties of the joints are far better than fusion-welded joints [250]. The paramount aim of this research is to identify the benefits and limitations of FSW while applying the same method to join different grades of steel.

Comparing with fusion welding methods, this FSW method has many advantages like higher mechanical properties, corrosion resistance property, lesser heat-affected zone, and FSW does not require any filler metals or special environments to join steel [251]. FSW showed a substantial saving in power consumption. The power consumed in fusion welding was four times or higher than FSW for fabricating similar joints. Also, FSW exhibits better microhardness in the weld zone compared to fusion welding. Moreover, the HAZ in FSW was narrower than that in fusion welding. Nevertheless, the fusion welding process produces higher amounts of harmful gases such as carbon monoxide and carbon dioxide to the surroundings (2.7 ppm and 346 ppm, respectively). At the same time, the FSW is a green method, which does not produce any harmful gases. [252, 253].

The review on the friction stir welding on steel shows that more research is needed to conduct similar and dissimilar steels and other alloys using hybrid and conventional FSW to reduce the joining costs and enrich the commercialization of friction stir welding in all-metal joining areas. The above review shows that most researchers concentrate on enhancing the strength of the joints, and a few are concentrated on the economic feasibility of the research. Numerous researches were conducted on similar stainless steel joints using PCBN like expensive tools; still, the strength gained was inadequate [254]. Using such tools are not recommendable due to its high expense. In dissimilar metal joining, the microstructure generated in the bimetal region is satisfactory, and the strength gained is insufficient to adopt this method to industry [255]. While laser-assisted FSW and induction heated FSW is widely used for welding steel and its alloys. One of the limitations of IA-FSW is that it can only use for ferromagnetic materials, while ultrasonic vibration-assisted friction stir welding can be adopted for any metals. High strength metals require high heat input to increase the metal flow to obtain a good joint. This volumetric material flow can be enhanced by clubbing an ultrasonic vibration unit with the friction stir welding unit. The significant advantage of ultrasonic vibration-assisted friction stir welding (UVAFSW) is that this ultrasonic vibration is directly given to the tool itself, so there is no energy loss. But since no evident research is undertaken on ultrasonic vibration-assisted FSW for steel and its alloys, this research area requires more attention.

While considering metal matrix composites (MMC), these MMC’s are poorly explored using friction stir welding. Dissimilar MMC joining by friction stir welding is still in the nascent stage. These limitations can be addressed by using hybrid friction stir welding and advance FSW tools. Still, enough researches are conducted on the cryogenic post-weld treatments for friction stir weld specimens. The study on pre and post-weld treatments’ effect on steels is a prerequisite in friction stir welding research. The impact of different profiled bobbin tools on steel friction stir welding needs more studies. These are the current research voids found in the friction stir welding of steels, which are required to address well.