On Novel Copper Based Alloys Development via Friction Stir Alloying

: Friction stir alloying (FSA) of commercially pure Cu with Ni, Zn, and Mg is implemented in the current study. Mechanical and microstructural aspects of the successfully fabricated alloy structure have been examined. Energy dispersive X-ray spectroscopy revealed a uniform distribution of alloying elements and coalescence at the atomic level. The compositional and grain size heterogeneity is managed in the stir zone, allowing for microstructural control with FSA. Thus, the present study is essential for the development of novel materials whose fabrication requires temperature well below the melting point of base metals. The alloying process is found to be accompanied by ultra-reﬁned grains, with the smallest grain size being ~0.44 µ m. The fabricated alloy managed to retain the FCC phase, and no brittle intermetallic compounds formed, according to X-ray diffraction. The fabricated alloy exhibits maximum and average microhardness enhancements of 18.4% and 6%, respectively. Tensile properties have also been investigated and correlated with microstructural morphology. A shift toward grain bimodality has also been documented, which is a highly sought-after property nowadays, especially to overcome the strength-ductility trade-off.


Introduction
Copper (Cu) and its alloys are widely used due to their excellent thermal and electrical conductivity. Cu alloys, which are the optimal option for today's energy and environmental problems, are extensively employed in nuclear fusion reactors at locations with exceedingly high heat flux [1][2][3]. Cu-based shape memory alloys (SMAs) and high entropy alloys (HEAs) form a group of functional materials that expand beyond Cu-based alloy's traditional electro-mechanical applications. Cu-containing HEAs have been discovered to possess antibacterial properties, as well as being corrosion resistant to marine biofilms [4]. Micro-electromechanical, electrical, and biomedical sciences are also investigating the use of Cu with varying microstructural morphology [5][6][7].
Friction stir alloying (FSA) is a relatively new development of the remarkable friction stir processing (FSP) technology, in which the matrix and reinforcement are both ductile metallic phases. The ideation of FSA also arises from the work of Shiri et al. [8], who performed friction stir welding by inserting the foils of Cu, Zn, Brass, and Cu-Zn-Ni between the abutting surfaces of Al plates and found the secondary materials to be diffused into the base metal. Chemical complexities and related mechanical adversaries that occur in fusion alloying processes can be easily avoided in FSA because alloying takes place at approximately half the melting point of the matrix material. Furthermore, FSA's ability to accommodate a wide variety of material heterogeneity allows it the modern-day alchemy for developing new alloy systems and libraries.
Shukla et al. [9] utilized FSA to scrutinize the phase transformations in Fe 40 Mn 20 Co 20 Cr 15 Si 5 (at. %) HEA after gradient Cu addition. In another study by Karthik et al. [10], the friction stir selectively alloyed Al stir zone with ∼2 wt. % Cu responded well to aging treatment, which emboldened the foundations of FSA. The HCP to BCC transformation due to the introduction of V in γ -FCC-based HEA via FSA has lately been investigated by Agarwal et al. [11]. AZ31 Mg alloy's mechanical capabilities have also been improved by varying Al additions using FSA [12]. In this case, the groove width was varied to modify the composition of Al in the stir zone. Several FSP studies consisting of metals, as both the matrix and reinforcement materials have also reported the formation of intermetallic compounds [13,14]. The discussed process has also been fruitful in stimulating transformationinduced plasticity (TRIP) and precipitates hardening by the addition of Ti in a twinning induced plasticity (TWIP) HEA [15]. Furthermore, FSA's ability to simultaneously produce ultra-refined grains while alloying is largely unmatched by other alloying processes. In another study, Li et al. [16] analyzed the consequences of Yttrium (Y) and Zirconium (Zr) co-alloying on microstructural transformation and the mechanical characteristics of the as-casted and the FSPed Mg-6Zn alloys. The yield strength, ultimate tensile strength, and elongation of the FSPed Mg-6Zn-1Y-0.5Zr alloys were significantly increased due to grain refinement, dispersion enhancement, and the customized composition. Mg-CuO and Mg-Cu-based surface nanocomposites were accomplished by Farghadani et al. [17] through FSP on the surface of AZ91 Mg casting alloy. The hardness of AZ91/Cu and AZ91/CuO nanocomposites was substantially enhanced, as well as AZ91/CuO nanocomposite displayed leading tensile strength and wear resistance. Similarly, FSP strengthened the hardness of the fabricated composite (Zirconia dispersed in Cu Matrix) as compared to Cu alone [18]. The primary reason for the increased hardness was the reduction in grain size and consistent distribution of zirconia.
Earlier research has demonstrated that FSP can be used to enhance the mechanical properties of an alloy by altering its microstructural characteristics. The grain size has a major impact on the strength of the alloy [19]. According to the Hall-Petch equation, there is an inverse relationship between the strength of the alloy and the grain size [20]. In FSP, the material is heated up due to friction and extreme plastic deformation, thus transforming the coarse-grained alloy into a fine-grained crystalline material. Naik et al. [21] implemented FSP to upgrade the mechanical and wear properties of Cu-Zr (Cu-0.18 wt.%Zr) alloy without compromising the electrical conductivity. They noticed a reduction in the average grain size from 40.5 to 4.6 µm and an improvement in hardness from 70 to 99 HV with a higher tool traverse speed. Among the different zones, the stir zone (SZ) has a massive influence on the microstructure and the base material's properties [22]. The volume of the base material processed during single-pass FSP is determined by the size and shape of the SZ. During FSP, two types of SZ shapes are commonly developed: basin and elliptical. According to Ma et al. [23], the shape of SZ is primarily governed by the heat input, which can be modified in FSP by changing the combination of tool rotation and travel speeds. Hsu et al. [24] used FSP to refine the SZ grains to 1-2 µm, resulting in an ultrafine-grained Al-Al 2 Cu with a high Young's modulus and strong compressive properties. Similarly, a fine equiaxed re-crystallized grain structure with an average size of 2.1 µm was formed in the SZ after FSP while dispersing graphene nano-platelets [25]. As a consequence, the nanocomposite's hardness, elastic modulus, and yield strength all increased dramatically. The FSP of Cu-Ni (70/30) by Mukherjee and Ghosh [26] resulted in reduced porosity, refined grains, higher hardness, minimized ductility, and increased corrosion, in contrast to laser-assisted direct metal deposition (DMD). Xie et al. [27] used a novel FSP aided laser cladding process to eliminate metallurgical defects. The coarse networks of laser cladding coating were converted into scattered nanoparticles, and the crack in laser cladding Ni-Cr-Fe coating was removed. The top surface grain size approached 300 nm, and hardness increased to over 400 HV. The investigation by Sabbaghian et al. [28] who fabricated a Cu-based composite by dispersing TiC through FSP demonstrated the many advantages of FSP yet again. They concluded that FSP formed a fine grain microstructure with a Crystals 2021, 11, 498 3 of 12 uniform particle distribution on the surface. They also illustrated that FSP strengthened the specimen's microhardness and wear resistance as a result of grain refinement and the presence of TiC particles in the FSP region. Priyadharshini et al. [29] explored the impact of tool traverse speed (of FSP) on the microstructural evolution, microhardness, tensile behavior, and tribological characteristics of CuNi/B 4 C surface composite. FSP of specimens at speeds greater than 30 mm/min resulted in undesirable microhardness, tensile strength, and wear resistance, according to the microstructural study. They found a maximum microhardness of 190 HV, a maximum ultimate tensile strength of 295 MPa, and a minimum wear rate of 210 × 10 −5 mm 3 /min at a traverse speed of 30 mm/min. According to Zangabad et al. [30], the experimental outcome on the fatigue properties of Al-matrix nanocomposites prepared by the FSP resulted in an increase in the fatigue strength. FSP is capable of preventing several defects, porosities, as well as improving the ductility and tensile strength of numerous alloys [31]. As a result, FSA which is a modern development of the, FSP must have the same characteristics, but further research is required. Although FSA's success has firmly been manifested in the above literature, the vast multi-dimensional elemental compositional space that is accessible through FSA remains chiefly unexplored. Studies concerned with the FSA of several extensively used metals, such as Cu, Zn, Fe, Ti, Mg, Ni, etc., are either very rare or utterly absent in the published works. As a result, the aim of this work is to use FSP to alloy four of the most commonly used metals, namely Cu, Zn, Mg, and Ni. This research will also look into the effect of the material's melting point on FSA. Hence, the alloying elements are chosen to cover a wide range of melting points, such as Zn melts at 420 • C, Mg melts at 650 • C, Cu melts at 1085 • C, and Ni melts at 1455 • C. To prevent the development of brittle intermetallic compounds, no more than four elements are considered for FSA. Cu and Ni would have been softened inadequately if Zn or Mg had been used as base materials. Ni as a base material, on the other hand, would lead to high softening temperature, that would provide formation enthalpy for intermetallic compounds. Thus, Cu is selected as the base material in the current research. Furthermore, surface alloying through FSA advances the current emphasis on surface behaviors and modeling [32]. The nature of FSA allows for microstructural control, both from the elemental and morphological perspectives. The alloying of Cu with Ni aids its corrosion resistance, solidus, and liquidus temperature, whereas the addition of Mg enhances ductility and workability. Further, Cu alloyed with Zn results in better creep strength and wear resistance. Therefore, alloying of Cu simultaneously with Mg, Ni and Zn is attempted in the present work through FSA. The simultaneous alloying with multiple metals using FSA is also the novelty of the present work. The characteristic zones, namely the SZ, thermo-mechanically affected zone (TMAZ) and the heat affected zone (HAZ) are examined in light of alloying for the first time. The ultra-refinement and transformed morphology of grains due to severe plastic deformation are scrutinized. The present research thus holds the potential to aid the development of meta-materials and unprecedented alloy systems through the FSA process.

Materials and Methods
Rolled commercially pure Cu plate (obtained from Tremor Alloys, Mumbai, India) of dimensions 180 mm × 60 mm × 4 mm was used for alloying. The chemical composition of the base metal is presented in Table 1. A robust vertical milling machine was utilized to perform FSA with the assistance of an in-house designed tool adapter and work fixture. A square slot with a side of 3 mm was milled on the Cu plate and filled with a mixture of pure Ni (50 wt. %), Zn (20 wt. %), and Mg (30 wt. %) powders. The higher weight proportion of elements with a relatively low melting point created voids/defects in the stir region during the trial experiments. Therefore, Ni (highest melting point) was used in the highest proportion, and Zinc (lowest melting point) was used in the smallest proportion among the alloying elements. Friction stir alloying also benefits from the fact that it is a solid-state process. As a result, keeping the foremost melting element in the smallest proportion ensures minimal melting during the alloying process. The average particle size of Ni, Zn, and Mg powders was~60 µm,~44 µm, and~100 µm, respectively. The slot was first covered by a pin-less tool having 20 mm shoulder diameter and then processed with a tool having shoulder, pin bottom, and pin root diameters of 23 mm, 10 mm, and 9.3 mm, respectively. A tapered cam-tri-flute profile was machined on the pin, having a length of 3.15 mm. High-speed steel (HSS) was used for tool material. Figure 1 shows a schematic diagram of the FSA experimental setup. The parameters indicated in the Figure 1, i.e., the tool rotational speed and tool traverse speed were set at 710 rpm and 50 mm/min, respectively. These parameters were established via optimized design using rigorous experimentation. smallest proportion ensures minimal melting during the alloying process. The average particle size of Ni, Zn, and Mg powders was ~60 μm, ~44 μm, and ~100 μm, respectively. The slot was first covered by a pin-less tool having 20 mm shoulder diameter and then processed with a tool having shoulder, pin bottom, and pin root diameters of 23 mm, 10 mm, and 9.3 mm, respectively. A tapered cam-tri-flute profile was machined on the pin, having a length of 3.15 mm. High-speed steel (HSS) was used for tool material. Figure  1 shows a schematic diagram of the FSA experimental setup. The parameters indicated in the Figure 1, i.e., the tool rotational speed and tool traverse speed were set at 710 rpm and 50 mm/min, respectively. These parameters were established via optimized design using rigorous experimentation. The most difficult aspect of FSP is achieving a uniform distribution of added materials, while simultaneously preventing agglomeration, voids, and surface craters. The tool traverse and rotational speed have the greatest impact on the above results. As a result, the above parameter setting was preceded by over 20 experiments in which one or more of the aforementioned issues were encountered. A standard metallographic sample preparation procedure was followed to examine the transverse cross-sections of the fabricated alloy. Before the examination, the samples were etched with a 1:1 mixture of HNO3 and H2O for 30 s, followed by rinsing with distilled water. Optical microscopy (OM) (QS Metrology, Delhi, India), Scanning Electron Microscopy (SEM) (Zeiss, Germany) and Energy Dispersive X-ray Spectroscopy (EDS) (Zeiss, Germany) were utilized to observe the morphology and composition of the developed microstructure. X-ray diffraction (XRD) (Rigaku, Tokyo, Japan) was utilized to scrutinize the crystal structure of the alloyed sample. Wire electric discharge machining was used to cut tensile samples, which were tested with a Tensometer (Kudale Instruments Pvt Ltd, Pune, India) of 20 kN capacity. Vickers microhardness was measured by a tester (Mitutoyo, Kanagawa, Japan) at a test force of 100 gm and a dwell time of 15 s, with an indentation separation of 0.5 mm. The microhardness of the base metal was 110 HV. The most difficult aspect of FSP is achieving a uniform distribution of added materials, while simultaneously preventing agglomeration, voids, and surface craters. The tool traverse and rotational speed have the greatest impact on the above results. As a result, the above parameter setting was preceded by over 20 experiments in which one or more of the aforementioned issues were encountered. A standard metallographic sample preparation procedure was followed to examine the transverse cross-sections of the fabricated alloy. Before the examination, the samples were etched with a 1:1 mixture of HNO 3 and H 2 O for 30 s, followed by rinsing with distilled water. Optical microscopy (OM) (QS Metrology, Delhi, India), Scanning Electron Microscopy (SEM) (Zeiss, Germany) and Energy Dispersive X-ray Spectroscopy (EDS) (Zeiss, Germany) were utilized to observe the morphology and composition of the developed microstructure. X-ray diffraction (XRD) (Rigaku, Tokyo, Japan) was utilized to scrutinize the crystal structure of the alloyed sample. Wire electric discharge machining was used to cut tensile samples, which were tested with a Tensometer (Kudale Instruments Pvt Ltd, Pune, India) of 20 kN capacity. Vickers microhardness was measured by a tester (Mitutoyo, Kanagawa, Japan) at a test force of 100 gm and a dwell time of 15 s, with an indentation separation of 0.5 mm. The microhardness of the base metal was 110 HV.

Macrostructure
The macrostructure of the as-fabricated friction stir alloy is depicted in Figure 2. The stirred region shows no sign of agglomeration of Ni, Zn, or Mg particulate. Instead, the SZ is macro-structurally consistent and continuous, which proves that alloying has been realized effectively. The TMAZ is sharper at the advancing side (AS), in comparison to the retreating side (RS) because of a greater degree of plastic deformation at the AS. It is noticed that the 9 mm 2 area of the slot was stirred by the pin having a peripheral area of 30 mm 2 and developed an SZ of 50 mm 2 size. The smaller expanse of the SZ at the bottom half of the alloyed region is widened at the top because of the mechanical stirring effect induced by the tool shoulder. A minute wormhole was discovered at the bottom of the SZ, presumably due to a material deficiency caused by the surface flash during processing.

Macrostructure
The macrostructure of the as-fabricated friction stir alloy is depicted in Figure 2. stirred region shows no sign of agglomeration of Ni, Zn, or Mg particulate. Instead, SZ is macro-structurally consistent and continuous, which proves that alloying has b realized effectively. The TMAZ is sharper at the advancing side (AS), in comparison to retreating side (RS) because of a greater degree of plastic deformation at the AS. It is ticed that the 9 mm 2 area of the slot was stirred by the pin having a peripheral area o mm 2 and developed an SZ of 50 mm 2 size. The smaller expanse of the SZ at the bott half of the alloyed region is widened at the top because of the mechanical stirring ef induced by the tool shoulder. A minute wormhole was discovered at the bottom of SZ, presumably due to a material deficiency caused by the surface flash during p cessing.

Microstructure
The scanning electron micrographs of the simultaneously ultra-refined and allo structure of the stirred region at two locations are shown in Figures 3a and 2b. The E mapping of Figure 3a is illustrated by Figure 3a1-a4. In contrast to the frequently obser agglomeration of non-metallic reinforcement in the matrix of composites fabricated FSP, the EDS mapping in Figure 3a1-a4 shows intimate intermixing and uniform dispers of Mg, Ni, and Zn.

Microstructure
The scanning electron micrographs of the simultaneously ultra-refined and alloyed structure of the stirred region at two locations are shown in Figures 2b and 3a. The EDS mapping of Figure 3a is illustrated by Figure 3a 1-a 4 . In contrast to the frequently observed agglomeration of non-metallic reinforcement in the matrix of composites fabricated by FSP, the EDS mapping in Figure 3a 1-a 4 shows intimate intermixing and uniform dispersion of Mg, Ni, and Zn.

Macrostructure
The macrostructure of the as-fabricated friction stir alloy is depicted in Figure 2. The stirred region shows no sign of agglomeration of Ni, Zn, or Mg particulate. Instead, the SZ is macro-structurally consistent and continuous, which proves that alloying has been realized effectively. The TMAZ is sharper at the advancing side (AS), in comparison to the retreating side (RS) because of a greater degree of plastic deformation at the AS. It is noticed that the 9 mm 2 area of the slot was stirred by the pin having a peripheral area of 30 mm 2 and developed an SZ of 50 mm 2 size. The smaller expanse of the SZ at the bottom half of the alloyed region is widened at the top because of the mechanical stirring effect induced by the tool shoulder. A minute wormhole was discovered at the bottom of the SZ, presumably due to a material deficiency caused by the surface flash during processing.

Microstructure
The scanning electron micrographs of the simultaneously ultra-refined and alloyed structure of the stirred region at two locations are shown in Figures 3a and 2b. The EDS mapping of Figure 3a is illustrated by Figure 3a1-a4. In contrast to the frequently observed agglomeration of non-metallic reinforcement in the matrix of composites fabricated by FSP, the EDS mapping in Figure 3a1-a4 shows intimate intermixing and uniform dispersion of Mg, Ni, and Zn.  The average grain size in Figure 3a,b is measured to be 2.63 µm and 2.1 µm, respectively. Therefore, ultra-refinement during alloying has also been reached throughout the SZ. Figure 4a,b represent the grain size distribution corresponding to Figure 3a,b, respectively. Figure 4b shows that the recrystallized grains at certain locations in the alloyed region have a tendency toward bimodality. The alloyed region has grains ranging in size from 0.44 µm to 7.5 µm.  The average grain size in Figure 3a,b is measured to be 2.63 μm and 2.1 μm, respectively. Therefore, ultra-refinement during alloying has also been reached throughout the SZ. Figure 4a,b represent the grain size distribution corresponding to Figure 3a,b, respectively. Figure 4b shows that the recrystallized grains at certain locations in the alloyed region have a tendency toward bimodality. The alloyed region has grains ranging in size from 0.44 μm to 7.5 μm.  An XRD analysis was performed in order to identify any possible phase transformations or the development of brittle intermetallic compounds in the fabricated alloy. Figure 6 shows the XRD results of the alloyed sample. The diffraction peaks seen at 2Ө values of 43.33°, 50.54°, 74.24°, and 90°, are corresponding to (1 1 1), (2 0 0), (2 2 0) and (3  An XRD analysis was performed in order to identify any possible phase transformations or the development of brittle intermetallic compounds in the fabricated alloy. Figure 6 shows the XRD results of the alloyed sample. The diffraction peaks seen at 2Өvalues of 43.33 • , 50.54 • , 74.24 • , and 90 • , are corresponding to (1 1 1), (2 0 0), (2 2 0) and (3 1 1) planes in pure FCC Cu. This indicates that during alloying, phase transformations and the development of brittle intermetallic did not occur. Furthermore, as evidenced by EDS findings, the fabricated alloy retains the FCC phase due to a high atomic proportion of Cu (greater than 90%). An XRD analysis was performed in order to identify any possible phase transformations or the development of brittle intermetallic compounds in the fabricated alloy. Figure 6 shows the XRD results of the alloyed sample. The diffraction peaks seen at 2Ө values of 43.33°, 50.54°, 74.24°, and 90°, are corresponding to (1 1 1), (2 0 0), (2 2 0) and (3 1 1) planes in pure FCC Cu. This indicates that during alloying, phase transformations and the development of brittle intermetallic did not occur. Furthermore, as evidenced by EDS findings, the fabricated alloy retains the FCC phase due to a high atomic proportion of Cu (greater than 90%). The light microscopy images taken at different locations on the alloyed macrostructure are elucidated in Figure 7. Figure 7b shows the shoulder affected SZ on AS clearly demonstrating the ultra-refinement of grains from HAZ to SZ. The material intermixing The light microscopy images taken at different locations on the alloyed macrostructure are elucidated in Figure 7. Figure 7b shows the shoulder affected SZ on AS clearly demonstrating the ultra-refinement of grains from HAZ to SZ. The material intermixing occurs at the RS SZ/TMAZ interface, as evident from Figure 7c. Owing to the induced strain and thermal cycle in the TMAZ, partial recrystallization occurs, which leads to larger grains as compared to the SZ (Figure 7d). The layer-by-layer deposition of material undergoing varying strain generates visible intercalations in microstructure, as shown in Figure 7e,f. The wide transition of TMAZ to HAZ on the RS can be observed in Figure 7g.

Property Evolution across Different Zones
The distribution of microhardness along the transverse section (as marked in Figure 1) of the alloyed specimen is represented by Figure 8. The ultra-refined alloy achieved a maximum microhardness of 130.2 HV, a peak enhancement of 18.4 % in comparison to the base. The augmentation in hardness is due to alloying and ultra-refinement, however, the role of grain size strengthening is greater.
The average hardness of the alloyed region is 6% higher than the base metal. Also, the alloyed region exhibits significant variation in microhardness across the transverse direction. The presence of numerous domains having disparate grain sizes and elemental compositions is the primary reason for this variation. The coarsening of grains in the HAZ due to prevalent heat results in the softening of this region. A comparison between the results of the current study and those in the literature is laid out in Table 2.
occurs at the RS SZ/TMAZ interface, as evident from Figure 7c. Owing to the induced strain and thermal cycle in the TMAZ, partial recrystallization occurs, which leads to larger grains as compared to the SZ (Figure 7d). The layer-by-layer deposition of material undergoing varying strain generates visible intercalations in microstructure, as shown in figures 7e and 7f. The wide transition of TMAZ to HAZ on the RS can be observed in Figure 7g.

Property Evolution Across Different Zones
The distribution of microhardness along the transverse section (as marked in Figure  1) of the alloyed specimen is represented by Figure 8. The ultra-refined alloy achieved a maximum microhardness of 130.2 HV, a peak enhancement of 18.4 % in comparison to the base. The augmentation in hardness is due to alloying and ultra-refinement, however, the role of grain size strengthening is greater.

Property Evolution Across Different Zones
The distribution of microhardness along the transverse section (as marked in Figure  1) of the alloyed specimen is represented by Figure 8. The ultra-refined alloy achieved a maximum microhardness of 130.2 HV, a peak enhancement of 18.4 % in comparison to the base. The augmentation in hardness is due to alloying and ultra-refinement, however, the role of grain size strengthening is greater. The tensile properties of base metal and test sample cut from the processed plate are given in Table 3. The wormholes in the SZ were machined out to measure the actual strength of the processed specimen. The fractographic SEM images corresponding to the failed Cu base metal and processed Cu tensile specimen are shown in Figure 9a-f, respectively. Since the SZ has been strengthened more than the base metal and the HAZ is weakened, morphological discontinuity exists at the TMAZ/HAZ interface. Therefore, the failure occurs from this location.    The partial recrystallization and possible diffusion of Ni, Zn, and Mg in the TMAZ result in a slight increase in the ultimate tensile strength (UTS) from 236 MPa to 240 MPa. However, unlike the TMAZ, HAZ consists of invariability in grain size, which results in a minor decrease in ductility. This is because, as stress increases during deformation in materials with a broad range of grain sizes, the larger grains make plastic deformation, while the smaller grains bear the load gradually, preventing sudden brittle failure. The fractured surface of the processed specimen has fewer dimples and ridges than the base metal, indicating that the specimen's ductility has decreased due to the uniform grain size of HAZ.

Conclusions
Alloying of Cu with Ni, Zn, and Mg is effectively achieved in the present study using the FSA process. The microstructural and mechanical properties of the fabricated alloy have been investigated in relation to the thermomechanical conditions that existed during the alloying process. The following conclusions can be drawn from this research: • FSA has been successfully used to fabricate alloys containing elements with a wide range of melting temperatures. The alloying elements (Ni, Mg, and Zn) are uniformly distributed in the stirred base metal (Cu) during FSA, but slight elemental heterogeneity is preserved across minute domains in the fabricated alloy.

•
The surface alloys manufactured through FSA show bimodal grain size distribution along with the ultra-grain refinement of the base metal. In the present study, the grain size of the fabricated alloy has reduced tremendously, ranging from 0.44 µm to 7.5 µm.

•
Owing to the disparity in grain size accompanied by the heterogeneous introduction of alloying elements, the average hardness of friction stir alloyed Cu is 6% higher than the base metal. FSA results in a peak hardness enhancement of 18.4 %, with the maximum hardness, obtained being 130.2 HV.

•
Strengthening of the fabricated alloy in the stir zone due to reduction in grain size and alloying results in a morphological discontinuity across the TMAZ/HAZ interface. If the whole substrate is subjected to tensile stress, the tensile failure occurs at this interface. However, the obtained UTS of 240 MPa in the current scenario is slightly greater than that of the base metal (236 MPa).