Trans. Nonferrous Met. Soc. China 30(2020) 2369-2380

Local inhomogeneity of mechanical properties in stir zone of friction stir welded AA1050 aluminum alloy

Xiao-chao LIU^1,2,3, Yun-qian ZHEN¹, Yu-feng SUN⁴, Zhi-kang SHEN^1,2,3, Hai-yan CHEN^1,2,3, Wei GUO¹, Wen-ya LI^1,2,3

1. School of Materials Science and Engineering, Northwestern Polytechnical University, Xi’an 710072, China;

2. State Key Laboratory of Solidification Processing, Northwestern Polytechnical University, Xi’an 710072, China;

3. Shaanxi Key Laboratory of Friction Welding Technologies, Xi’an 710072, China;

4. School of Materials Science and Engineering, Zhengzhou University, Zhengzhou 450001, China

Received 31 December 2019; accepted 4 June 2020

Abstract: The local inhomogeneity of the stir zone in friction stir welded face-centered cubic metal was investigated, which has multiple activated slip systems during plastic deformation, by selecting commercial AA1050 aluminum alloy as an ideal experimental material. The local inhomogeneity was evaluated by uniaxial tensile tests using small samples with a 1 mm gauge length. The corresponding microstructural parameters such as grain size, misorientation angle distribution, and micro-texture, were quantified by the backscattered electron diffraction technique. A comprehensive model was used to reveal the microstructure-mechanical property relationship. The experimental results showed that the uniaxial tensile property changes significantly across the weld. The maximum ultimate tensile strength (UTS) occurred in the center of the stir zone, which was 99.0 MPa. The weakest regions were located at the two sides of the stir zone. The largest difference value in UTS reached 14.9 MPa, accounting for 15% of the maximum UTS. The analysis on the structure-mechanical property relationship suggests that the micro-texture change with the location formed during the rotational material flow is the main reason for the local inhomogeneity.

Key words: friction stir welding; face-centered cubic metal; local inhomogeneity; mechanical properties; micro-texture

1 Introduction

Fine-grained metals and their alloys have been fabricated by many different techniques in the last decades [1,2]. However, their weldability had always been poor until the friction stir welding (FSW) technique was invented [3-8]. It has been generally accepted that the FSW is a thermo-mechanical coupling solid-state joining process [9-11]. Compared with conventional fusion welding, the heat input in the FSW is much lower [12-14]. However, to further reduce the heat input in the FSW of fine-grained materials, smaller sizes of the tool and thinner workpieces [4,5] or larger loads combined with lower rotating speeds [15,16] should be adopted. Rapid cooling friction stir welding (RCFSW) is another promising method to obtain fine grains under conventional welding conditions. In RCFSW, the cooling medium is jetted on the tool and the workpiece. In this way, the heat effect can be reduced to the minimum, and even eliminated [17-21]. Therefore, the stir zone (SZ) is forced to be the key factor for the whole joint quality. In this background, much attention should be paid to the local inhomogeneity of the SZ.

The inhomogeneity in the mechanical properties of the FSW joint has been widely concerned. Much attention was paid to the close-packed hexagonal (HCP) metals subjected to the FSW, such as magnesium, titanium, and their alloys [22-26]. PARK et al [22] reported that the micro-texture with {0001} plane parallel to the surface of the tool pin was formed in the FSW of AZ61 Mg alloy. Under the influence of this micro-texture, the joints fractured at the transition zone between the SZ and the base metal (BM) during the uniaxial tensile tests perpendicular to the welding direction (WD). WANG et al [24] pointed out that the {0001} micro-texture was formed by the basal slip during the simple shear deformation because the basal slip is most easily to be activated during the deformation of Mg alloys. For the titanium and its alloys, owing to the fact that the prismatic slip is more easily to be activated than the basal slip, the  simple shear texture was formed in the SZ of titanium [25]. LIU et al [26] found that the  texture gave rise to the inhomogeneous deformation during the tensile deformation of pure titanium FSW joint. In summary, for the HCP metals, the inhomogeneity in the mechanical properties of FSW joints is mainly caused by the strong texture formed due to the activated single slip system during the material flow.

However, for the face-centered cubic (FCC) and body-centered cubic (BCC) metals (Al, Cu, Fe, etc.), which have multiple activated slip systems during the plastic deformation, the previous studies mainly focused on the inhomogeneity caused by different strains, thermal history or dissimilar materials in the joint of FSW [27-39] or friction stir spot welding [40,41]. For example, the SZ, the thermo-mechanically affected zone (TMAZ), and the heat affected zone (HAZ) have always different mechanical properties for various materials [27-34]. SIMAR et al [27] investigated the local and global mechanical properties of the FSW joint of aluminum alloy 6005A-T6. It was found that although the hardness was similar in the SZ and the HAZ, the strain hardening capacity of the weld center was larger than that of the HAZ, and thus fracture occurred in the HAZ of the joint. SHEN  et al [29] reported that the fracture location varied with the welding speed in the FSW of copper. The average hardness in the SZ first decreased and then increased with the increase in welding speed, but it was nearly constant in other regions within the joint, and thus the weakest region in the joint varied.

The dissimilar materials joining also gives rise to inhomogeneous mechanical properties within the joint [35-41]. DONATUS et al [35] investigated the variation in elemental compositions across the weld zones of a dissimilar friction stir weld between AA5083-O and AA6082-T6 alloys to illuminate the inhomogeneous microhardness distribution in the SZ and the TMAZ. KANGAZIAN and SHAMANIAN [36] investigated the mechanical behavior of the dissimilar FSW joints between the SAF 2507 super duplex stainless steel and the Incoloy 825 Ni-based superalloy and found that the joint owned the similar strength to the Incoloy 825 parent metal.

In summary, little attention was paid to the inhomogeneous mechanical properties in the SZ of FCC and BCC metals with multiple activated slip systems during the plastic deformation. In this study, a simple FCC metal, aluminum alloy, was used as an ideal experimental material to investigate the inhomogeneity of the SZ in the FSW of FCC metal.

2 Experimental

The cold-rolled plates of commercial AA1050 with dimensions of 300 mm × 60 mm × 3 mm were butt-welded by RCFSW using liquid carbon dioxide as a cooling medium. A FSW tool containing a concave shoulder of 12 mm in diameter and a smooth cylindrical probe of 2.8 mm in length and 4 mm in diameter was used. To avoid the kissing bond and tunnel defects, the welding speed and rotating speed were 200 mm/min and 1200 r/min, respectively. The welding temperature was estimated to be ~0.7T_m (T_m is the melting point) [42,43]. The tool tilt angle was 2.5°. After welding, the welds were cut to small non-standard tensile samples by spark cutting. The dimensions of the tensile samples are shown in Fig. 1(a). Such small-size tensile specimens were also used in the determination of the joint quality of laser beam welded and diffusion bonded joints [44-46]. To highlight the local inhomogeneity, the parallel length of the tensile samples is only 1 mm and the width is 1.7 mm, about 10 times the revolutionary pitch. The samples were cut every 1 mm from the weld center to the advancing side (AS) and the retreating side (RS) and marked as “Center, AS-1 to AS-5 and RS-1 to RS-5”, respectively. The samples were first thinned from the weld bottom to ~1.2 mm in thickness and then thinned from the top surface to ~1 mm in thickness by heavy-grade sandpaper. Finally, the samples were polished by finer-grade sandpaper to a very good surface finish. The entire process was finished in flowing water. The tensile tests were conducted at room temperature using an Instron-type testing machine at a constant cross- head speed of 3.3×10^-4 mm/s.

Fig. 1 Relative locations of small tensile samples to weld and their dimensions (a) and locations of EBSD measurements on weld cross-section (b) (unit: mm)

The microstructural parameters such as grain size, misorientation angle distribution, and micro- texture at different locations corresponding to the small tensile samples were characterized and quantified by the electron back-scattered diffraction (EBSD) technique. For the EBSD measurement, the specimens were firstly ground by sandpaper in anhydrous ethanol, and then mechanically polished by using Al₂O₃ polishing paste to the mirror finish. At last, the electrolytic polishing with a solution of V(HClO₄):V(C₂H₅OH)=1:4 at 20 V and 0 °C was conducted. The EBSD measurements were carried out on a JEM-7001F field-emission scanning electron microscope (FE-SEM) operated at 15 kV using the TSL OIM^TM system. The measured positions were along the line with 0.5 mm from the top surface and arranged every 1 mm from the center to the AS and the RS on the weld transverse cross-section, as shown in Fig. 1(b). The step size of the EBSD scanning was 0.3 μm. More than 5000 grains (enclosed by boundaries with misorientation angle higher than 2°) were calculated by EBSD software automatically to obtain the microstructural parameters. A 15° criterion was used to differentiate the low angle boundary (LAB) and high angle boundary (HAB). The Vickers hardness of the corresponding positions was measured with a load of 50 g and a dwell time of 15 s.

3 Results

3.1 Tensile properties

Figure 2 shows the uniaxial tensile test results. The stress-displacement curves of different samples scatter to a great extent. The sample “Center” has the highest ultimate tensile strength (UTS) and the tensile displacement. From the center to the AS and the RS, the UTS and tensile displacement show a trend of gradual decline. At the same time, the samples with the same distance to the weld center have similar stress-displacement curves.

The detailed yield stress (σ_0.2) and UTS of the samples are listed in Table 1. The yield stress ranges from 62.8 MPa (RS-5) to 71.9 MPa (RS-1). The maximum difference is 9.1 MPa, which is ~12.7% of the maximum yield stress. The UTS ranges from 84.1 MPa (RS-5) to 99.0 MPa (Center). The maximum difference reaches 14.9 MPa, which is 15% of the maximum UTS. These differences values indicate that the local inhomogeneity in the mechanical properties of the SZ should be concerned.

Fig. 2 Tensile engineering stress curves at different locations of weld

Table 1 Yield stress and ultimate tensile strength at different locations of weld

3.2 Microstructure

To understand the above inhomogeneous mechanical properties in the SZ, the microstructures at different locations were characterized by the EBSD technique. Figure 3 shows the grain microstructures at the locations from AS-5 to Center and then to RS-5. The black and white lines in the EBSD inverse pole figure (IPF) maps denote the HABs and LABs, respectively. Different colors stand for different grain orientations under the specimen coordinate system. From Fig. 3(a) to  Fig. 3(k), the dominant hue changes gradually with the locations from AS-5 to RS-5. This means that the micro-textures vary greatly at different  locations. This phenomenon has been widely reported in previous studies [47-49]. The grain size, HAB and LAB fractions, and misorientation angle distribution from AS-5 to RS-5 were summarized and shown in Fig. 4. The average grain size is nearly constant at different locations, i.e. ~2.5 μm. The number fraction of LAB presents a gradually decreasing trend from AS-5 to RS-5. At AS-5, due to the mixing with the TMAZ (Fig. 3(a)), the LAB fraction is significantly higher than that at the other locations. The lowest LAB fraction occurs at RS-5, which is only 19%. All the misorientation angle distributions at different locations present a two-peak distribution with one peak at 2° and the other peak at ~50°. Such a bimodal misorientation distribution is like that of aluminum alloy subjected to severe cold deformation and subsequent annealing [50,51]. This indicated that the full recrystallization occurred in the SZ during the RCFSW [52,53].

Figure 5 shows the (111) pole figures at the locations from AS-5 to RS-5. In the pole figure, the vertical direction stands for the normal direction (ND) of the workpiece, and the horizontal direction is in accord with the transverse direction (TD) of the workpiece. The pole figures vary regularly from AS-5 to RS-5. This should be attributed to the rotational material flow during the FSW [54-56], which results in different shear directions and shear planes at different locations.

The detailed micro-texture changes from AS-5 to RS-5 are shown in Fig. 6. As shown in Fig. 6(a), the original pole figures can be divided into several single-components. The orientation of a single-component can be read from the orientation distribution function (ODF). The number fraction of every single-component was calculated using a tolerant misorientation angle of 15° by the EBSD analysis software. Four main components, their number fractions and Taylor factors are summarized and listed in Table 2. The remnant part was considered as a random orientation. Their distributions in the space of (0 ≤ φ₁, Φ, φ₂ ≤ π/2) are shown in Fig. 6(b). The three-dimensional coordinate position of the colorful symbol denotes the Bunge Euler angles of the corresponding single-component. The area of the symbol stands for its relative fraction. Two orientations dominated the micro-texture at a specific location. Both rotated from the RS to the AS, and their intensity increased gradually, as shown by the dotted arrows in Fig. 6(b). It should be emphasized that some rotations cannot be seen intuitively only in the space of (0 ≤ φ₁, Φ, φ₂ ≤ π/2) due to the multi- symmetry of FCC crystal. These micro-textures are known as  simple shear texture after rotating around the TD axis [57]. These rotational micro- textures are formed by rotational material flow. This is a typical phenomenon occurring in the FSW.

Fig. 3 EBSD IPF maps showing grain microstructures at different locations of stir zone

Fig. 4 Changes of microstructure parameters at locations from AS-5 to RS-5

Fig. 5 (111) pole figures at locations from AS-5 to RS-5

4 Discussion

4.1 Strengthening mechanisms

To reveal the relationship between the local mechanical properties and the corresponding microstructures, the following discussion tries to use a comprehensive model to calculate the contribution of different strengthening mechanisms.

Fig. 6 Micro-texture changes from As-5 to RS-5

For polycrystalline aluminum alloy, it has been generally accepted that the boundary strengthening and dislocation strengthening are the main strengthening mechanisms [50,51] and the orientation factor M cannot be neglected especially for textured samples. The yield stress σ_y or flow stress σ_f is usually described by the following formula [50]:

σ_y(or σ_f)= σ_gb+σ_dis+Mτ_tot                                (1)

where σ_gb and σ_dis represent the contributions of grain boundary strengthening and dislocation strengthening, respectively. The orientation factor M is often referred to as the Taylor factor which is significantly dependent on the crystalline texture and the direction of the tensile axis. τ_tot is the critical resolved shear stress (CRSS), which consists of intrinsic CRSS, solid solution strengthening, and precipitation strengthening. For commercial AA1050, the CRSS mainly refers to the intrinsic CRSS due to its low alloy elements. In this study, the CRSS of fully annealed AA1050 is taken as τ_tot and it is about ~9 MPa [58,59]. σ_gb can be described by Hall-Petch relationship [60]:

                              (2)

where k is a constant and d_gb is the average grain size. KAMIKAWA et al [50] suggested that the LABs with misorientation angles above 2° act as conventional grain boundaries in terms of strength contribution and thus the parameter d_gb in Hall-Petch relationship is the average value of the grain and subgrain sizes (i.e. enclosed by boundaries with misorientation angles above 2°). The remnant boundaries with misorientation angles below 2° and grain interior dislocations are considered to contribute to the dislocation strengthening. σ_dis can be described by the following formula [50]:

                         (3)

α is a constant (0.24) [50,61], G is the shear modulus (26 GPa for aluminum), b is the Burgers vector component (0.286 nm for aluminum), and ρ_dis is the dislocation density. In this study, owing to the full recrystallization occurring in the SZ, the dislocation density in the grain interior was very low [62].

4.2 Grain boundary strengthening

The slope k of Hall-Petch relationship for grain boundaries strengthening at a given strain can be obtained from the relationship between the hardness and the square root of the grain size [62].

Table 2 Four main single-components of micro-texture, their number fractions and Taylor factors (along with TD) at corresponding locations from AS-5 to RS-5

Generally, the slope k is obtained by fitting the linear relationship between the yield stress and the square root of the grain size. However, in this study, as the yield stress test is easily influenced by the local micro-texture, the hardness was used to calculate the slope k. In the hardness test, the applied stress is in multiple directions and thus the influence of micro-texture can be ignored. As shown in Fig. 7, the circle symbols stand for the data obtained in the RCFSW of AA1050 in this study and our previous study [52]. The square symbols denote the data obtained in the conventional FSW of AA1050 [62]. Through fitting the data, the slope k is ~24.8 HV/μm^-1/2 in the RCFSW of AA1050. This value can be converted to ~81 MPa/μm^-1/2 using the equation H≈3σ, where H is the diamond pyramid hardness and σ is the stress at a strain of 8% during a tensile test [63]. This equation is workable when the material is without appreciable work hardening [62,64]. The slope k in the RCFSW is a little higher than that of the conventional FSW. SATO et al [62] suggested that a higher value of k corresponds to the higher dislocation density. This indicates that the dislocation density in the RCFSW is a little higher than that in conventional FSW. This is because the rapid cooling suppresses the annealing during the cooling stage [52,55,65]. It is worth mentioning  that when the slope k is calculated by using the hardness, the contribution of the dislocation strengthening has been included in the slope k. Therefore, dislocation strengthening should not be considered separately again.

Fig. 7 Hall-Petch relationship between microhardness and square root of grain size in RCFSW of AA1050

In this study, the measured hardness at different locations and the corresponding tensile stress at 8% strain are summarized and listed in Table 3. They roughly match the equation H≈3σ. The small deviation between them may be caused by the higher dislocation density in the RCFSW. These results indicate that the above approximate treatment of which the slope k in the Hall-Petch relationship is converted to ~81 MPa/μm^-1/2 from ~24.8 HV/μm^-1/2 is acceptable.

Table 3 Comparison between diamond pyramid hardness (H) and tensile stress (σ) at strain of 8%

4.3 CRSS strengthening

For the orientation factor M, the classical Taylor model is used to calculate the values of M at different locations, in which five slip systems are as summed to be active. For texture free FCC metals, the Taylor factor M_tf=3.07 [66]. As for the present textured samples, M can be calculated by the following equation:

                (4)

where f_i is the number fraction of the grain with the same orientation, M_i is the corresponding Taylor factor. The remnant fraction of the grains is assumed as random orientation and the corresponding Taylor factor is M_tf. M_i is dominated by the specific orientation (φ₁, Φ, φ₂) and the tensile stress axes. For FCC metal, the detailed calculation of M_i refers to Ref. [66]. In this study, M was calculated by the EBSD analysis software   directly. The calculated results have been listed in Table 2.

4.4 Comparison of flow stress

As the equation H≈3σ works when the strain is 8%, the following discussion will focus on the comparison between the calculated flow stress and the experimental value at the strain of 8%.

Figure 8 shows the comparison of the calculated flow stresses at the strain of 8% with the experimental values at different locations. Here, the experimental flow stress is expressed by the engineering stress with a strain of 8%. This is reasonable because the true tress is roughly equal to the engineering stress at such a small strain. For comparison, the corresponding UTS and the microhardness of the SZ are also presented in Fig. 8. Notably, the microhardness in the SZ is nearly uniform. However, the UTS and the flow stress present an obvious unimodal distribution. The stress values in the SZ center are significantly higher than those on the two sides. The variation trends of the calculated and experimental flow stresses match each other very well from AS-5 to RS-5 with a maximum error of ~2.3 MPa (~ 3%).

Fig. 8 Comparison between calculated flow stress and experimental value at 8% strain

Figure 9 shows the contributions of the   grain boundary strengthening and the CRSS strengthening. According to the results, the grain boundary strengthening is the main strengthening mechanism for fine-grain aluminum produced by the RCFSW. However, the grain boundary strengthening at different locations is almost the same. The difference in the flow stress is mainly caused by CRSS strengthening. This indicates that the variation of the orientation factor M caused by the rotational micro-texture plays a very important role in the local inhomogeneity of the SZ. Therefore, it can be inferred that for any fine-grain FCC metals fabricated by FSW, the larger the CRSS is, the moreremarkable the local inhomogeneity of the SZ is.

Fig. 9 Contributions of grain boundary strengthening and CRSS strengthening to calculated flow stress and experimental flow stress

The phenomenon that the microhardness at different locations in the SZ is nearly the same is a typical feature in the fine-grain aluminum alloy weld fabricated by the FSW [4,5]. According to the microstructure investigation, the grain size and misorientation angle distribution also have no significant difference at different locations of the SZ. However, the local inhomogeneity of the mechanical properties in the SZ is nonnegligible. Taking the UTS as an example, the highest value (in the center of the SZ) is 15% larger than the lowest value (on the two sides of the SZ). This phenomenon should be concerned when fabricating fine-grain FCC materials by the FSW. Also, it is worth mentioning that the local inhomogeneity of mechanical properties in the SZ of the FSW is greatly different from that in the fine-grain materials fabricated by other deformation processing methods, such as equal channel angular pressing and high-pressure torsion [59,67]. The latter two arise from the locally non-uniform plastic deformation and are reflected in the hardness, grain size, and dislocation density. However, in the FSW, the local inhomogeneity is mainly caused by rotational micro-texture. Therefore, if the micro- texture intensity in the SZ can be weakened, the local inhomogeneity will be reduced and even eliminated. MIRONOV et al [42] found that lower rotating speed is beneficial for the weakening of the micro-texture intensity in the SZ. LIU et al [52] suggested that rapid cooling can suppress the micro-texture sharpening during the cooling stage. These methods can be considered to relieve the local inhomogeneity of mechanical properties in the SZ caused by the rotational micro-texture.

5 Conclusions

(1) The tensile strength of friction stir welded AA1050 in the center of the SZ is significantly higher than that on the two sides of the SZ. The maximum difference is 14.9 MPa, accounting for 15% of the local maximum UTS.

(2) The grain size and misorientation angle distribution have no obvious change at different locations. However, the micro-texture presents an obvious rotational trend around the tool axis (normal direction of the workpieces).

(3) The flow stress at a strain of 8% during the tensile tests is successfully predicated by taking the grain boundary strengthening and the CRSS strengthening into consideration, which agrees well with the experimental value.

(4) The local inhomogeneity of mechanical properties in the SZ originates from the CRSS strengthening, which is determined by the rotational micro-texture at different locations to the tensile direction.

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铝合金AA1050搅拌摩擦焊搅拌区的力学性能局部不均匀性

刘小超^1,2,3，甄云乾¹，孙玉峰⁴，申志康^1,2,3，陈海燕^1,2,3，郭 伟¹，李文亚^1,2,3

1. 西北工业大学 材料学院，西安 710072；

2. 西北工业大学 凝固技术国家重点实验室，西安 710072；

3. 陕西省摩擦焊接工程技术重点实验室，西安 710072；

4. 郑州大学 材料科学与工程学院，郑州 450001

摘  要：以铝合金AA1050为例，研究在塑性变形过程中具有多重滑移系的面心立方金属搅拌摩擦焊搅拌区的力学性能局部不均匀性。采用标距为1 mm的小试样单轴拉伸测试评估力学性能的局部不均匀性；通过电子背散射衍射技术对搅拌区内局部的晶粒尺寸、取向差角分布和微织构等结构参数进行量化；利用综合分析模型揭示显微组织和力学性能的关系。实验结果表明，局部的单轴拉伸性能沿着焊缝宽度方向发生显著变化。抗拉强度的最大值出现在搅拌区中心，达到99.0 MPa；最小值位于搅拌区的两侧。抗拉强度之间的最大差值达14.9 MPa，为抗拉强度最大值的15%。基于结构-力学性能关系的分析可知，材料旋转流动过程中形成的随位置变化的微织构是导致局部力学性能不均匀性的主要因素。

关键词：搅拌摩擦焊；面心立方金属；局部不均匀性；力学性能；微织构

(Edited by Wei-ping CHEN)

Foundation item: Project (51905437) supported by the National Natural Science Foundation of China; Project (2019M653726) supported by the China Postdoctoral Science Foundation; Project (3102019QD0407) supported by the Fundamental Research Funds for the Central Universities, China

Corresponding author: Zhi-kang SHEN; Tel: +86-29-88495240; E-mail: zhikangshen@nwpu.edu.cn

DOI: 10.1016/S1003-6326(20)65385-7