1 Introduction

Friction stir welding (FSW) is finding many industrial applications, slowly replacing traditional fusion welding technologies. The process is now firmly established for aluminum alloys, which still account for most FSW applications [1]. However, as argued by Brassington et al. [2], new possible applications involving titanium alloys can be conceived in the space industry due to their unique characteristics (i.e., high strength-to-weight ratio, excellent corrosion resistance, and compatibility with mostly propellants). Ti6Al4V is among the titanium alloys whose friction stir weldability has been studied the most. As reported by Mironov et al. [3], friction stir welding of this titanium alloy presents several issues, such as tool wear, the need for an inert gas shield to minimize oxidation, a cooling system to avoid overheating the tool and spindle system, and a welding machine capable of sustaining very high axial forces (i.e., three times those required for aluminum alloys). An overview of friction stir weldability of Ti6Al4V was proposed by Edwards and Ramulu [4], ranging between 65 and 100 mm/min for the welding speed and 170 and 300 rpm for rotational speeds. Additionally, in a follow-up study, they measured temperature in the nugget zone (NZ) while welding Ti6Al4V sheets between 1000 and 1200 \(^{\circ }\)C above the \(\beta \)-transus temperature (980 \(^{\circ }\)C) [5].

Concerning Ti6Al4V microstructure induced by friction stir welding, Kitamura et al. [6] found an equiaxed \(\alpha \) and elongated \(\beta \) grains microstructure in NZ when the temperature is below the transus point. Instead, when the temperature in the NZ overcomes the \(\beta \)-transus temperature, the microstructure is characterized by a fully \(\beta \) transformed microstructure with basket-weave lamellar \(\alpha \)/\(\beta \) structure, as demonstrated by Su et al. [7]. Wu et al. [8] focused on the final microstructure obtained in the thermo-mechanically affected zone (TMAZ). According to the authors, a complex process characterizes the final microstructure. First, the dislocation generated by the deformation arranges, forming subgrain boundaries due to dynamic recovery. Then the high temperature induces the DRX by transforming low-angle boundaries into high-angle boundaries generating newly refined grains. In the end, new dislocations are introduced in the refined recrystallized grains. Instead, the weakest zone of the joint is the heat affected zone (HAZ), as stated by Fall et al. [9]. The authors found a modified microstructure compared to the base material, with a bimodal structure with lamellar and equiaxed \(\alpha \) grains justified by temperature reached over the \(\beta \)-transus temperature.

Research over the past decade has provided an overview of FSW joints of Ti6Al4V, at least in traditional butt-butt configuration. However, complex geometries, such as T-welds, are essential for structures in transportation, such as light ships and aircraft fuselages [10]. As reported by Tavares et al. [11], different types of T-welds can be realized. The main configurations are the following:

  • T-lap joint, with two plates, one horizontal (skin) and one vertical (stringer) (Fig. 1a);

  • T-butt joint with complete penetration, i.e., horizontal sheets lateral surfaces faying on the vertical one in the middle (Fig. 1b);

  • T-butt joint without complete penetration, i.e., horizontal sheets lateral surfaces faying between them (Fig. 1c);

Fig. 1
figure 1

Principal friction stir welding T-joint configurations

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Tavares et al. [11] showed the excellent mechanical properties of the T-butt joints with the stringer between the two horizontal slabs (Fig. 1b). Whereas Zhao et al. [12] highlighted the critical issue in T-joint configuration in the support’s radius fillet and the tool geometry’s choice. Both aspects play a role in kissing bond defects at the radius between the skin and the stringer, affecting the joint’s mechanical strength. Derazkola et al. [13] pointed out the importance of the tilt angle in friction stir welding in T configuration, identifying the optimal value to maximize mechanical properties at 2\(^{\circ }\). However, while we are beginning to have relevant information on friction stir welding T-joint aluminum alloys, more research still needs to be conducted on T-welds of titanium alloys. The authors acknowledge two research groups working on this issue in the past years, whose first papers were published in 2020. Su et al. [14] first studied the weldability of Ti-4Al-0.005B alloy in T-butt joint configuration with complete penetration performing two welds along the two interfaces obtaining a free defects joint. In their second work, Su et al. [15] compared the mechanical properties of joints obtained, welding either the two interfaces or once, centering the tool in the mid-thickness of the vertical slab. The authors achieved sound welds with both strategies reaching good tensile strength while the elongation was always around 2% against the 20% reached in the BM. The welding strategy affected the fracture position with crack occurring at the corner for single-weld and at the HAZ for double-weld T-joints. Higher mechanical properties were found with the single-weld strategy. The second research group, Campanella et al. [16], began work instead on T-lap welding of the Ti6Al4V alloy. In the first work, they demonstrated the feasibility T-lap joint with this alloy through friction stir welding, which until then had always been welded by melting-based welding technologies. In their subsequent work [17], the authors proposed a numerical model to simulate the final microstructure of the T-joint on the same test already carried out in the previous work.

In this investigation, for the first time, Ti6Al4V T-joints with complete penetration are manufactured to assess their friction stir weldability and mechanical properties. Two tools and different combinations of welding parameters are employed to study their influence on the joint quality (i.e., defectiveness and mechanical strength). Thus, the goal is to understand the microstructure and mechanical strength of the joints and how their weldability is affected by the welding conditions and tools, improving the state-of-the-art in FSW in more complex geometry welds of Ti6Al4V alloy to identify new potential applications.

2 Experimental procedures

2.1 Friction stir welding

Three parts friction stir welding T-joints with complete penetration are performed on FSW dedicated machine MTS ISTIR PDS. The welds 140 mm long along the rolling direction are executed on Ti6Al4V. The horizontal sheets are \(140 \times 55 \times 1.2\) mm (L x W x T), while the vertical is \(140 \times 21.2 \times 1.2\) mm.

The alloy chemical composition and mechanical properties obtained characterizing base material with three tensile tests and five hardness HV\(_{0.3}\) indentations are listed in Table 1.

Table 1 Ti6Al4V chemical composition according to the supplier certificate and mechanical properties defined through testing
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Figure 2a displays the base material microstructure, consisting of \(\beta \) phases placed along \(\alpha \) phases grain boundaries. Two types of frustum-shaped cobalt-based alloy tools are used [18]. Pins are featureless, with tapered profiles that are cylindrical and triangular with rounded edges, while the shoulder is concave with an angle of 5\(^{\circ }\). Three welding conditions are tested for both tools carrying out two welding for each configuration.

Figure 2b displays the tools, their geometrical information, and the employed welding parameters.

Fig. 2
figure 2

Information about base material, tools and welding parameters

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Rounded corner fillets in the fixture characterize mild steel support to ensure good formability of the corners in T-joints. The radius of rounded edges is fixed at 1 mm. A ZrO2 coating is applied on the support surface to avoid sticking between titanium alloy sheets and steel support due to the local high temperature possibly reached during the process. Additionally, argon protects the welding zone during the stirring to avoid welding seam oxidation. Figure 3 displays the experimental setup showing the whole configuration.

Fig. 3
figure 3

Experimental setup details

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2.2 Welds characterization

The welds’ strength is characterized by tensile tests along the skin and Vickers hardness measurements. An Instron electromechanical testing machine is used for uniaxial tensile tests equipped with a 100-kN load cell and cross-head speed of 1 mm/min at room temperature. Subsize specimens are tested according to ASTM E8M [19]. On each T-joint, three tensile samples and one for hardness and microstructure analysis are obtained through water-jet cutting, as displayed in Fig. 4. Five tensile tests for each configuration are performed, and deformation is measured with digital image correlation (DIC) through the post-processing software Aramis. A Zwick ZHU 2,5 durometer with a load of 2.94 N and a duration time of 10 s is used for hardness measurements. The indentation pattern in the weld cross-section is displayed in Fig. 4 and is performed on advancing and retreating sides. After polishing and etching with Kroll’s reagent, the same samples are used for microstructural observation.

Fig. 4
figure 4

Indentations pattern and samples water-jet cutting scheme (Base material = BM, heat affected zone = HAZ, thermomechanically affected zone = TMAZ, shoulder affected zone = SAZ, nugget zone = NZ)

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The etched samples are inspected using an optical microscope (Olympus PMG3) for microstructural detail and a digital microscope (Keyence VHX6000) for macroscopic observations. One specimen’s microstructure is characterized using a Zeiss Ultra Plus FESEM microscope equipped with an X-Max Oxford Instruments system and x-ray silicon drift detector (SDD) of 20 mm\(^2\) to acquire high magnification micrographs (5000X to 25000X) and to perform semi-quantitative analysis of the chemical composition by energy dispersive spectroscopy (EDS). For tool wear analysis, 3D profiles of tools (new and after welding) and weld exit holes are rebuilt with the Alicona Infinite Focus 5 G. Tools’ profiles are analyzed with a laser scanner (Riftek model RF627) to compare the volume of the as-manufactured and worn tools automatically.

3 Results

The feasibility of Ti6Al4V friction stir T-welds is presented in the following subsections: First, tool wear is studied to provide insight into tool wear evolution and preferable sites using two Co-based tools. Afterwards, the weld quality is assessed by analyzing the macrographies to identify the weldability window. Mechanical properties and fractography of the samples are discussed to establish the strength achievable when friction stir welding Ti6Al4V complex geometries. In the end, the microstructure of one selected sample is analyzed to evaluate microstructural modifications in the different zone to infer the maximum temperatures reached.

3.1 Tool wear

Among problems linked to friction stir welding of titanium alloys, tool wear is one of the most critical due to severe and rapid tool wear, as reported by Wang et al. [20]. Materials able to resist the thermomechanical stresses developed during FSW of titanium are often costly, difficult to machine, and wear rapidly [21]. The tools used in this work are made of Co-based alloy. They are easily machinable and could represent a cost-effective FSW tool material solution for high-strength alloys such as titanium [22]. However, there is little information regarding the wear of this type of tool during the FSW of titanium alloys. Park et al. [23] analyzed cobalt-based tool wear during friction stir welding of Ti6Al4V alloy. According to the authors, these tools are cheap, easily castable, and machinable while standing good mechanical properties over 1000 \(^{\circ }\)C. At the same time, the wear rate is high when welding titanium alloys due to intermittent sticking on the tool surfaces, causing adhesion wear. Recently, Du et al. [24] characterized the wear mechanisms of a Co-based tool when friction stir welding TA5 alloy. The flat pin tip mainly wear for abrasion, while at the pin root and shoulder, wear is governed by adhesion and diffusion.

The qualitative analysis of tool wear is given by following the evolution of the weld exit-hole to infer tool wear by comparing the trace left by the pin when retreated between welds. Change in the tool shape is then revealed by the mark left by the tool. Figure 5a compares a new tool A and the same after 720 mm welds.

Fig. 5
figure 5

Details of tool and exit-hole to explain the tool wear qualitative analysis

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From the top view, the complete transformation of pin features is remarkable. The surface of the pin tip is significantly reduced after the welds. Additionally, the tapered pin lateral surface profile is difficult to recognize due to wear with its shape approaching a straight cylinder. The important parameter for the next steps is \(h_o\), represented in Fig. 5b. This height represents the distance between the pin tip and the tangent to the shoulder concavity. The \(h_o\) parameter is selected because it is possible to find the same analyzing the tool exit-hole profiles, as displayed in Fig. 5c and d. By rebuilding the 3D profiles of the mark left by the tool after retreating, including the hole and the shoulder-workpiece interaction zone, it is possible to obtain the parameter \(h_{WC}\) (Fig. 5d). It represents the same quantity as \(h_o\) after the distances traveled by the tool. The height \(h_{WC}\) is calculated for several tool exit-holes to follow the wear, and its evolution for both tools is displayed in Fig. 5e. While the height \(h_o\) for both manufactured tools is around 1.25 mm, it increases fast and continuously with the millimeters traveled, reaching 1.6 mm after 720 mm of welds, growing more than 30%. Considering that the pin’s tip is less sensitive to wear, as visible in Fig. 5b, the increased height is mainly justified by tool wear in the inner part of the shoulder. To locally analyze the wear, the 3D volume of the new and worn tools are compared. The overlapped volume and details on wear in the pin and shoulder are displayed in Fig. 6.

Fig. 6
figure 6

Local analysis of tool wear: details on pin and shoulder

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The wear drastically changes between shoulder and pin. In the shoulder, the maximum height difference measured between the new and worn tool ranges between 0.5 and 0.6 mm. For the pin, both the bottom and lateral surfaces should be considered. On the bottom surface, the wear is low in the center, with values below 0.1 mm, and it increases in the outer part reaching the edges. Additionally, while for tool A (Fig. 6a), the lateral wear of the pin is significant, i.e., 0.15 mm, the same does not apply for tool B (Fig. 6b). The difference is due to the different pin profiles. The rounded trigonal edges characterizing the manufactured tool A disappear during welding due to wear and explain the 0.15 mm wear. Contrarily, the lateral surface of the tapered cylindrical pin B is less sensitive to wear, not being characterized by stress concentration points (i.e., edges). Wear of Co-based FSW tool when welding Ti6Al4V is prominent and quick, and the preferred site is the inner part of the shoulder. Pin features are mechanical and thermal stress concentration spots wearing quickly, completely modifying the initial tool shape.

3.2 Weld quality

Weld cross-sections are displayed in Fig. 7. Independently from the tool and the welding parameters, no flow-related defects are found. The thermo-mechanical condition led to a proper material flow, guaranteeing the correct filling of the cavity left by the tool while advancing.

Fig. 7
figure 7

T-joint cross-sections

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Hence, using a tilt angle of 2\(^{\circ }\), welding speed between 60 and 100 mm/min, and rotational speed between 150 and 200 rpm, sound Ti6Al4V T-joints are obtained. The weldability window is valid for two tool shapes, tapered cylindrical and trigonal with rounded edges.

The weld cross-sections’ features are shown in Fig. 8. Due to the severe tool wear occurring during each weld, black traces left in the stirred zone can be observed, as illustrated in Fig. 8, but they are present to a different extent in all the cross-sections (Fig. 7). The trace can be found at the interface shoulder-workpiece and the advancing side in the transition between SAZ and NZ. In some welding configurations, those traces can also be found in the nugget zone (Fig. 7b–d and f). EDS analysis within the NZ of sample A-WC2 is carried out to clarify the nature of the observed dark traces. Figure 9a shows the macro area of the EDS spot (red oval solid line), while Fig. 9b displays the exact point corresponding to the spectrum (red cross). Observing the element revealed by the analysis in Fig. 9c, despite the titanium and the two major alloying elements, i.e., aluminum and vanadium, four extra elements such as tungsten (2%), nickel (1.9%), cobalt (1.3%) and chromium (0.5%) appeared. All of them are contained within the Co-based alloy FSW tool according to the patent [18].

Fig. 8
figure 8

A-WC1 cross-section details: tool particles within the weld zone, kissing bond and thinning

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Hence, the four extra chemical components prove tool particles within welds. When analyzing cross-sections, the variations in the dark traces observed in the stirred zone are justified by the welding parameters combination leading to a different amount of tool wear and material flow. Another critical aspect to highlight is the thinning of the weld cross-section due to the shoulder forging action. The thinning gradually evolves from the weld seam’s limit to the center of the nugget zone. By measuring the thickness of the sheet in the middle point between those two extremes, as illustrated in Fig. 8, it is possible to estimate the thinning at 0.1 mm. Similar thinning results were observed for all the joints.

Concerning the weld corners, two aspects are worth noting. Firstly, in all joints, both corners are not perfectly smooth and rounded because of the slight lateral movement of the horizontal sheets during the process due to unsuitable lateral clamping. Secondly, the kissing bond defects are observed on both sides of the weld, as highlighted by the dotted and dashed red squares in Fig. 8. The kissing bond defect has been well addressed by Sato et al. [25], representing in friction stir butt welding a partial remnant of the unwelded butt surface below the stir zone attributed to the insufficient plunging of the welding tool during FSW. As assessed by Tavares et al. [11], these defects can occur in friction stir T-welds in the corners because of the insufficient pressure applied by the tool due to insufficient plunge depth or unsuitable combination of pin-radius fillets. In the T-joints presented, kissing bonds on both sides always occurred independently of welding parameters and tools. This result suggests an unsatisfying combination of pin geometry-support radius leading to partially unwelded surfaces. The main difference between the two sides is the notable thicker oxide layer on the retreating side. The asymmetry is explainable, with the tool shifting a few tenths of a millimeter towards the advancing side, as observable from the cross-section.

Fig. 9
figure 9

EDS of darker particles in the NZ: a details of the NZ, b SEM observations to highlight the spot and c spectrum

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Hence, while the main problem of internal voids is avoided, kissing bonds in the corners and cross-section thinning of the joints occur during welding. Also, tool particles are observed in the joint, starting from the interface shoulder-workpiece, and spreading within the nugget zone.

3.3 Mechanical properties

Vickers hardness measurements across the welds are shown in Fig. 10a and b, for tools A and B, respectively.

Fig. 10
figure 10

Vickers hardness profiles in the cross-section

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Analyzing the profiles, the low influence of welding configurations and tools on the hardness is worth noting. As reported by Kitamura et al. [6], the main factors affecting hardness in Ti6Al4V friction stir welds are the thermal cycle intended as cooling rate and peak temperature. The two elements are affected by welding and rotational speeds. In the welding configuration employed, the range of welding parameters is limited to minimize tool wear, maximizing the tool life. Therefore, the similarities in hardness are induced by very similar thermal cycles guaranteeing similar microstructures for all the welds in the different zones. The higher hardness values are found in the SZ, the narrow decreasing zone may be linked to the TMAZ, and in the HAZ, the hardness is slightly lower or equal to the BM. The temperature may explain the result with peak temperature in HAZ below the \(\beta \)-transus one and resulting microstructure similar to the BM. Additionally, the asymmetrical hardness in the heat affected zones is already reported by different authors [26, 27] and may be justified by material plastic deformation and peak temperature differences between AS and RS.

The mechanical strength of the T-joints are analyzed through engineering stress–strain curves estimating ultimate tensile strength (UTS), yield strength (Y\(_{0.2\%}\)), and elongation at break (\(\varepsilon _{b,\%}\)). The engineering stress is calculated using the area obtained by multiplying the base material thickness of 1.2 mm by the width depending on the water-jet cutting. Engineering stress–strain curves for the different T-joints, organized for tool and compared to the BM, are shown in Fig. 11 displaying just one out of the five tensile tests performed for each configuration.

Fig. 11
figure 11

Engineering stress–strain curves for the different welding configurations

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On the one hand, the influence of the welding configuration on the assembled mechanical properties is negligible, with the curves almost overlapping, as displayed in Fig. 11a and b. On the other hand, differences between curves in the two figures can be observed, with the joints performed with tool B reaching slightly higher maximum stresses and deformation. While UTS and Y seem to approach, at different extents, the base material, the elongation at break for all joints is low, without reaching 2%. Hence, T-joints can reach high UTS and Y values while breaking as soon as they enter the plastic regime, as illustrated in Fig. 11. The remarkable UTS and Y result from a proper material flow induced by the pin, avoiding internal defects and consequent premature failure of the samples. Contrarily, the surprisingly low deformation reached at failure may be explained through the flaws highlighted in the previous sections, such as the kissing bond at the corners and the cross-section thinning. Mechanical properties mean values and standard deviations of the T-joints according to the tool and welding configuration are summarized in Table 2. Considering the standard deviations, the negligible differences in mechanical properties reached by employing different welding parameters are worth noting. Joints executed with tool B are more stable than the ones obtained with tool A, in which WC2 maximizes the mechanical properties.

Table 2 Mechanical properties for T-joints manufactured with all wedling configurations
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The critical issue raised by analyzing T-joint strength is the low elongation at break reached in all configurations. The T-joints fractography may elucidate this aspect. Despite very similar mechanical properties and hardness, two fracture types characterized the T-joints. The ones obtained through WC1 and WC2, independently from the tool, all failed in the heat affected zone on the retreating side. Instead, only 50% of the B-WC3 follow the same failure path. In order to clarify the fracture paths, frames before and after fracture during tensile tests and the longitudinal strain field before fracture obtained through DIC are displayed in Fig. 12a–c. The strain is concentrated on the retreating side before fracture, as shown in Fig. 12b, and the failure occurs in the HAZ below the shoulder. The second fracture type occurred at the kissing bond on the retreating side. All T-joints welded in WC3 and with tool A failed at the kissing bond against the 50% of B-WC3. Frames before and after fracture and the strain field before the fracture are shown in Fig. 12d–f. In this case, the fracture occurs in the corner even if the longitudinal strain distribution is similar to Fig. 12b, i.e., concentrated in the HAZ on the RS. Hence, the joints’ behavior is similar when subjected to traction independently from the welding configuration and tool, but in WC3, the fracture path can change. Differences in the fracture behavior of the samples can be clarified by analyzing fracture surfaces through SEM observations. Figure 13a shows the fractured surfaces in specimen B-WC2, which failed between the shoulder affected zone (SAZ), i.e., the zone affected mainly by the shoulder, and the HAZ. Three distinct fracture surfaces can be distinguished, in agreement with [15] when studying the fracture behavior of Ti-4Al-0.005B. They can be divided into shear and flat fracture zones and are illustrated in Fig. 13a.

Fig. 12
figure 12

a–c Detail of fracture in the HAZ (B-WC3); d–f detail of fracture at the kissing bond on the RS (A-WC3)

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The flat zone includes zone A, corresponding to the crack initiation zone located in the upper part of the joint. Zone B, the crack propagation area, is also included in the flat surfaces and is shared between SAZ and HAZ. Zone C represents the shear zone and corresponds to the rapid fracture zone containing only the HAZ. Zone B and C are displayed in Fig. 13b and c, while details at higher magnification are illustrated in Fig. 13d and e. Both zones are characterized by deep dimples (red arrows), typical features of ductile fracture surfaces [28]. Dimple ridges and valleys between zone B and C are similar, while they increase in size in zone C. The size difference is justified by the microstructure characterizing the SAZ and the HAZ, with a smaller grain size in the SAZ due to the stirring and consequent grain refinement. Instead, ridges and valleys are similar because of the significant plastic deformation undergone by the two zones due to the strain localization. Macroscopic fracture surface of B-WC3 fractured in the kissing bond zone is shown in Fig. 14a. The same fracture zones can be identified, but the path is inverse, starting at the bottom and ending at the upper surface. The fracture begins at the kissing bond (zone A), develops in the center (zone B), and ends in the upper part of the weld (zone C). The fracture occurs within the NZ of the friction stir weld. Zone B and C are displayed in Fig. 14b–d, and e and are characterized by small, shallow dimples (red arrows). Differences in dimple size can be explained by the different microstructures in the fractured zone (SAZ/HAZ for B-WC2 and NZ for B-WC3). Qin et al. [29] demonstrated the proportionality between grain size and dimples diameters. It is known that the stirring pin and shoulder action in friction stir welding is responsible for a refined equiaxed microstructure in the stirred zone. Consequently, the smaller dimples found in B-WC3 fracture surfaces and zone B in B-WC2 (Fig. 13b and d) are explained through the microstructural differences in grain size between stirred zones (by the pin and shoulder) and heat affected zone.

Fig. 13
figure 13

Fracture in the HAZ (B-WC2): a Fractured sample; b and d crack propagation zone; c and e rapid fracture zones

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However, ridges and valleys are more pronounced, i.e., deeper dimples, in fractured surfaces from the specimen failed between SAZ and HAZ. The dimples depth difference is explained by the same elongation at break found and the similar strain field before fracture displayed in Fig. 12, but different fracture path and zone. For sample B-WC3, while during the test, the strain is localized in the weaker and thinner zone, i.e., the heat affected zone, due to the kissing bond defects, the crack starts suddenly to propagate at the kissing bond tip instead of occurring at the HAZ for necking and void coalescence in the dimples zone. Hence, the fracture in B-WC3 differs from that in B-WC2, with shallow dimples because of the strain localized in HAZ during the tensile test. In contrast, the fracture propagates at the kissing bond tip characterizing lower plastic deformation undergone by the failed zone. Deeper dimples in Fig. 13d and e due to the substantial amount of plastic deformation occurring before the final separation [30], against the small and shallow dimples found in Fig. 14d and c, confirm the hypothesis. Therefore, although there are differences in fracture surfaces between samples, elongation at break are similar due to the previous strain localization in the heat affected zone, then overcome in certain configuration by the sudden crack propagation at the kissing bond tip.

Fig. 14
figure 14

Fracture in the HAZ (B-WC3): a Fractured sample; b and d crack propagation zone; c and e rapid fracture zones

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3.4 Microstructure

To infer the microstructure in the different zone composing the Ti6Al4V T-joints, SEM observations in the base material (BM), heat affected zone (HAZ) and nugget zone (NZ) on sample A-WC2 are performed. Two observations at different magnifications for each zone are shown in Fig. 15. The base material is characterized by bimodal equiaxed \(\alpha \)-phase and intergranular \(\beta \)-phase, highlighted in Fig. 15d with red and black arrows, respectively.

Approximating the NZ, within the heat affected zone, the microstructure begins to transform due to the thermal cycle undergone by the base material. The impact of the rapid heating and subsequent cooling induced by friction stir welding on the advancing side (AS) is revealed in Fig. 15b and e. HAZ microstructure still resembles the BM, but some modifications occurred. Simultaneously, intergranular \(\beta \)-phases grow (black solid arrow), and intragranular \(\beta \)-phases (black dotted arrow) appear in the \(\alpha \) matrix due to the temperature reached during the process. Those observations suggest that the maximum temperature reached within the HAZ did not overcome the \(\beta \)-transus temperature. Otherwise, microstructure in the HAZ should have presented both \(\alpha / \beta \) grains and equiaxed-\(\alpha \)-phases, as reported by Li et al. [31]. However, the increased size of intergranular \(\beta \) and the higher density of intragranular \(\beta \) phases in \(\alpha \) grains suggested a temperature in the \(\alpha -\beta \) transition range, with a consequent percentage increase of \(\beta \) against \(\alpha \) phases. Additionally, the slight differences in values observed between the HAZ on the AS and the BM (Fig. 10) are justified by the growth of \(\beta \) intergranular phases in HAZ leading to the few points reduction observed.

Focusing on the NZ, as always occur in FSW, the microstructure is wholly transformed compared to other zones composing the joints due to the high temperature and strain rate reached during the stirring. Figure 15c and f display the transformed microstructure characterized by equiaxed (EA) \(\alpha \) and elongated (E) \(\beta \) grains. As discussed for the HAZ, the absence of the lamellar structure suggests that the peak temperature in the nugget zone never exceeded the \(\beta \)-transus temperature. The results of [6] reinforce this hypothesis. The stir zone obtained below the \(\beta \)-transus point is given by dynamic recrystallization achieved in friction stir welding, i.e., high temperature and strain rate in the stirred zone, generating the refined and equiaxed \(\alpha \) grains structure.

Fig. 15
figure 15

SEM observations in the different zones: a and d BM, b and e HAZ (AS), c and f NZ

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4 Discussion

Tool wear analysis proved a massive reduction in the shoulder height even after less than a meter of welding. However, worn tools did not negatively affect FSW joints’ defectivity and mechanical properties. The tools ensured the correct material mixing in all samples avoiding internal voids in the selected welding configuration. Also, static mechanical properties are not influenced by tool wear states. These positive results prove the reliability of FSW-worn tools, at least for the wear encountered in our test. Of course, the definition of critical limit over which the weld standards are not met (i.e., internal voids or decreasing mechanical properties) may be helpful in future works. Furthermore, wear should be quantified to comprehend the welding impact on the base material employed for our tools. For this purpose, their volumes before and after welding are quantified. The difference in volume divided by the welded distance provides the wear rate. In order to simplify the volume estimations, we assume axisymmetric wear, considering that the whole surface wears evenly. The pin of tool B is axisymmetric, while in A, it is not due to the trigonal external surface making the volume estimation not straightforward. However, most wear is localized on the shoulder. Hence, we neglect the pins’ surface to compare wear, bypassing the problem of non-axisymmetry of pin A. The areas considered for volume calculation are visually highlighted in red and orange in Fig. 16a. The wear rate is similar for both tools, considering the same base material, weld conditions, and millimeters traveled. Based on the bar graph displayed in Fig. 16b, the average shoulder wear rate between the two tools is around 0.045 mm\(^3\) per millimeter of welding when friction stir welding Ti6Al4V with a Co-based alloy tool. Such a high wear rate might be a financial, i.e., the unit cost of the tool, and a time-wise problem, i.e., frequent tool changes, if we imagine adopting FSW to weld titanium T-joints in mass production.

Fig. 16
figure 16

Information about shoulder wear for both tools

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Friction stir welding of T-joints with complete penetration obtained in the investigated combination of process parameters and tools does not present critical internal voids. Consequently, welding speeds between 60 and 100 mm/min and rotational speeds between 150 and 200 rpm are suitable for avoiding internal defects related to material flow. To quantify the mechanical strength of the joints, it is helpful to directly compare it to the base material through the efficiency parameter, i.e., the ratio between the joint and base material (BM) mechanical properties. Because of the similar values obtained with the different welding configurations and to get an overall view of joints’ strength, efficiencies are calculated by averaging the mean values for each welding configuration of UTS, Y, and \(\varepsilon _b\). Efficiencies are reported in Table 2. The tapered cylindrical pin (tool B) gives better mechanical properties with a 94% and 89% efficiency for UTS and Y, respectively, while the tapered trigonal rounded edges pin maximum efficiencies are 88% and 86%. Nevertheless, hardness measurements are very similar between all welding configurations, as displayed in Fig. 11, not corroborating mechanical testing results. Discrepancies in UTS and Y average values can be explained as they are calculated using the theoretical section of 1.2 mm of the base material. However, when observing cross-sections in Fig. 7, not uniform thicknesses are evident due to the arc pattern generated by the shoulder-workpiece interaction. Variations can therefore be linked to the differences between true and theoretical sections. Nonetheless, the efficiency of UTS and Y consistently above 80% using the conservative thickness without considering the thinning is a comforting result for the first comprehensive report on the mechanical properties of complex friction stir welds of Ti6Al4V. The efficiencies of elongation at break are very low for both tools staying below 20% of the base material. Samples failed either at the HAZ or the corner, always on the RS. The good mechanical properties, UTS and Y, comparable with the base material, are confirmed by microstructural observations in HAZ and NZ. The heat affected zone presents a similar microstructure as the base material, with the growth of \(\beta \) intergranular phases, while the NZ is characterized by refined equiaxed \(\alpha \) and elongated \(\beta \) grains. The microstructural observations justify the hardness measurements, with hardness values in HAZ very similar to BM and higher values reached in NZ due to the significantly smaller size of the grains, according to the Hall-Patch relation. The microstructural modifications induced by friction stir welding suggest that \(\beta \)-transus temperature has never been exceeded during the test, avoiding the typical lamellar structure reported by Liu et al. [32]. As demonstrated by Kitamura et al. [6], the best mechanical properties of Ti6Al4V are achieved when the stirred and heat affected zone temperatures are below the \(\beta \)-transus point, avoiding the formation of lamellar structure. Similar mechanical properties and hardness measurements are achieved in the three welding configurations. Hence, the weldability of friction stir T-joints below the \(\beta \) transition temperature has been assessed, and maximum mechanical properties can reach 94% and 89% of the base material for UTS and Y, respectively.

The critical issues are the low elongation at break, the strain concentration and fracture on the RS. Hardness measurements show negligible differences between RS and AS and cannot explain why the fracture always occurs on the RS. Low elongations at break below the 20% of the BM are reported by Su et al. [15]. Failures are similar to the ones we encountered, i.e., in the HAZ. According to the authors, the rupture in the HAZ started at the upper surface of the T-joints, characterized by an arc pattern considered a stress concentration feature likely becoming a point of crack initiation. The arc pattern can be observed in our welds in Fig. 7 and linked to the thinning of the section. Instead, Liu et al. [27] also encountered elongation at break comparable to our results in butt configuration when friction stir welding Ti6Al4V. The authors explained the early rupture through section thinning and the consequent yielding only in the zone below the shoulder and deformation not evenly distributed in the samples. Hence, as already reported by other authors, the joints mainly failed in the heat affected zone, below the shoulder-workpiece interaction zone, because of the section thinning. The smaller section in the center of the weld leads to strain concentration in this area instead of evenly distributing it in the whole sample length. Additionally, all the samples fractured at the retreating side in our tensile tests. To understand the strain concentration on the RS, as highlighted in Fig. 12b and e, it is noteworthy analyzing the upper surface of the joints in Fig. 17.

Fig. 17
figure 17

Cross-section analysis to explain the section reduction

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The shoulder forges the sheets’ upper surface, generating a dissymmetrical arc pattern in the cross-sections. The arc shape of the weld seam is linked to the tilt angle used, i.e., 2\(^{\circ }\), and the tool plunge depth. However, the non-symmetry is associated with the increased flash formation on the retreating side, detected in all welds. The excessive flash on RS is driven by non-optimal material flow below the shoulder. The shoulder can only partially keep the material beneath its concave surface resulting in material ejection on the RS rather than its containment and release to ensure a uniform and symmetrical surface between AS and RS. Additionally, tool particles (Co, W, Ni and Cr) mixed with the workpiece are visible in the stirred zone and deposited on the upper surface, mainly on the AS. These particles released due to tool wear cyclically at the end of each rotation could favour dissymmetries. Therefore, the excessive flash on the RS and the tool particles concentrated in the upper surface on the AS leads to different true sections between the two sides of the friction stir welds before starting the tensile tests. Consequently, a slightly smaller area at the RS might be the cause of the 100% rate of failure in the RS for joints breaking in the HAZ. Cross-section heights are measured at different points at AS and RS between the nugget zone and the shoulder diameter mark in the areas represented by the dashed double arrow in Fig. 17c to identify the minimum height on both sides. Measurements revealed a difference (\(\Delta h = h_{min,AS}~-~h_{min,RS}\)) ranging from 0.04 to 0.10 mm between AS and RS, confirming all configurations’ minimum cross-section on the retreating side.

To understand why some WC3 configurations failed in the corner, studying the evolution of kissing bond defects with welding parameters is noteworthy. Details of the kissing bond defects on the retreating side of the T-joint obtained with the trigonal pin are displayed in Fig. 18.

Fig. 18
figure 18

Kissing bond defects details on the retreating side of joints obtained with tool A

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It is remarkable how its extension and direction evolve with welding parameters. Based on the measurement illustrated in Fig. 18, on the one hand, WC3, i.e., highest welding and rotational speeds, leads to the longer and more inclined oxide layer, a possible critical nucleation site for crack propagation during tensile tests. On the other hand, the hottest configuration, WC2 (highest rotational speed and lowest welding speed), leads to a shorter and almost vertical oxide layer. Welding parameters affect the extension and direction of the kissing bond defect. Similar results are obtained with tool B (tapered trigonal rounded edges profile), with thicker oxide layers on the retreating side and welding parameters influencing its length and direction. WC3 gives the most critical kissing bond defect in length and orientation (i.e., towards 45\(^{\circ }\)). Hence, the joints’ behavior is similar when subjected to traction independently from the welding configuration and tool. Nevertheless, if the oxide layer is long, thick, and inclined towards traction, the fracture path can change with a sudden and instantaneous crack growth starting from the kissing bond defect on the RS until the upper surface. Differences in fracture behavior have been shown through SEM observations in the crack propagation and rapid growth zones. Dimples in samples fractured in the HAZ are big in diameter and deep, i.e., a significant distance between ridges and valleys, as shown in Fig. 13. Contrarily, dimples in the sample fracture in NZ (Fig. 14), with the crack starting to propagate at the kissing bond tip, are small and shallow. Size differences are linked to the microstructure in the fractured zones. At the same time, the ridges’ lower extension is caused by the less amount of plastic deformation undergone by the material [30] in the fractured zone, considering the previous strain concentration in the HAZ. This phenomenon explains why the sample presents similar elongation at break and strain field distribution before fracture, even if fracturing in two different zones.

In conclusion, in the welding configurations tested in this work, independently from process parameters and tools, all the joints presented cross-section thinning and kissing bonds while avoiding internal defects. The chosen tilt angle and plunge depth induce the thinning, independently of welding and rotational speeds. Despite the section thinning, kissing bonds have always occurred, and their extension and orientation are affected by process parameters. A longer pin could have reduced or completely removed the interface’s oxide layer. However, being aware of the vertical steep thermal gradient existing when friction stir welding Ti6Al4V [33], the longer pin could wear fast because of the more resistant stirring material at the bottom, soon leading to a shorter pin reencountering the same problem. Hence, finding the optimal FSW T-welding setup is undoubtedly challenging due to several factors affecting the final result. Besides kissing bond and section thinning we encountered in all configurations leading to poor elongation at break (\(\sim \)15% of BM), in the best case, joints’ yield strength and ultimate tensile strength reached 87% and 96% of the base material. These excellent mechanical properties suggest that by solving the problem of kissing bond and thinning problem, complex assemblies that are certainly competitive with melting-based welding technologies may be achieved, taking advantage of all the environmental and automation benefits of friction stir welding.

5 Conclusions

This work demonstrated friction stir weldability of Ti6Al4V T-joints with complete penetration. Based on the present results, the following conclusions are drawn.

  1. 1.

    The investigated weldability windows (rotational speed between 150 and 200 rpm and welding speed between 60 and 100 mm/min) avoided critical internal voids, independently from the pin's shapes. However, minor flaws, such as small kissing bonds at corners and section thinning, are observed in all joints.

  2. 2.

    The similar hardness distributions observed in the different configurations suggested a uniform microstructure in the investigated weldability windows. Based on microstructural observations, temperatures in the various zones are below the \(\beta \) transus point.

  3. 3.

    Tensile tests have demonstrated the outstanding strength of Ti6Al4V friction stir T-welds. They reached very high mechanical properties, with the best case efficiencies of 96% and 87% for UTS and Y, respectively. However, the critical problem of elongation at break not higher than 15% of the base material is established.

  4. 4.

    Samples always shown the strain concentration in the HAZ on the retreating side due to cross-section thinning leading to the early fracture in the HAZ. Instead, some joints with the highest rotational and welding speed are characterized by a different fracture path starting from the kissing bond defect because of a more pronounced kissing bond defect (i.e., longer and inclined towards the traction direction). Hence, both factors, pronounced kissing bonds and section-thinning on one side, are responsible for low elongation at break. They can be avoided by adopting the correct plunge depth and combination of pin and support radius fillet.

  5. 5.

    Co-based alloy tool wears in Ti6Al4V friction stir welding has proved to be a critical issue, with relevant amounts of Co, W, Ni and Cr found within the nugget zone. The wear rate is very high and estimated at around 0.045 mm\(^3\) per millimeter, with a reduction in the shoulder height of about 0.5 mm after 720 mm of welding in its inner part closer to the pin root. Nevertheless, apart from tool particles within the weld cross-section, tool modifications induced by wear did not affect either weld quality or mechanical properties.