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Cu-10Zn interfacial observations
Acta Materialia 166 (2019) 324e334
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Elucidation of interface joining mechanism during friction stir welding through Cu/Cu-10Zn interfacial observations Huihong Liu*, Kohsaku Ushioda, Hidetoshi Fujii Joining and Welding Research Institute, Osaka University, Ibaraki, Osaka, 567-0047, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 October 2018 Received in revised form 31 December 2018 Accepted 3 January 2019 Available online 4 January 2019
A butt friction stir welding (FSW) process was performed on Cu and Cu-10Zn plates. Microstructural evolution along the Cu/Cu-10Zn interface retained in the FSW joint was systematically investigated in order to clarify the interface joining mechanism during the FSW. The initial joining interface is completely unbonded and has lots of large oxides distributed on it. The initial large oxides are then fragmented into small particles and the oxide layers initially formed on the interface are ruptured, which produces fresh surfaces on both the Cu and Cu-10Zn sides. The fresh surfaces are then joined at an atomic level under a compressive force perpendicular to the interface, while the small oxide particles are still dispersed on the interface with voids formed around the particles. A strain-induced grain boundary migration driven by the stored energy difference associated with the dislocation density difference will occur, which assists to extinguish the voids retained beside the oxide particles. Finally, a sound interface joining without oxide particles and large voids remaining on the interface is obtained during the FSW. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Friction stir welding STEM Microstructure Grain boundary migration Interface joining mechanism
1. Introduction Friction stir welding (FSW) is an innovative joining technique, which was invented by The Welding Institute (TWI) of the UK in 1991 originally aiming to join Al and its alloys that are difficult to fusion weld [1]. In FSW, a non-consumable rotating tool with the designated shoulder and pin is inserted into the interface of two butting work-pieces. The heat is mainly generated by the friction between the tool and the work-pieces, and it softens the workpiece materials around the tool, making them plastically deformable. The rotating tool then traverses along the butting line of the work-pieces accompanied by transporting the surrounding materials from the leading side to the trailing side of the tool to produce a joint [2]. Since FSW enables a joining in the solid state, it can be adopted to avoid or minimize the fusion-welding-associated problems such as formation of coarse columnar grains, high residual stress, severe welding distortion, and solidification defects as well as macro-segregation, which inevitably degrade the quality of the fusion weld joints. Moreover, since a fine and equiaxed recrystallized microstructure can be obtained in the friction stir welds by severe plastic deformation and elevated temperature, and
* Corresponding author. E-mail address: [email protected] (H. Liu). https://doi.org/10.1016/j.actamat.2019.01.004 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
the formation of a heat-affected-zone, which is generally the fracture site in a fusion-welded joint, can be effectively suppressed by the relatively low welding temperature, sound weld joints with good mechanical properties can be produced by the FSW [3e5]. Nowadays, FSW, which is considered as a green and nextgeneration joining technology, has been extensively applied in low-melting-point materials, such as Al and Mg alloys [6e10], and has also been expanded to many other high-melting-point materials including Cu, Fe and Ti alloys due to the development of highly durable tools [11e18]. Most of the FSW researches have focused on clarifying the physical fundamentals of FSW such as heat transfer, material flow, strain and strain rate, and microstructural evolution during FSW. Nandan et al. [19] investigated the three-dimensional heat transfer and material flow during FSW of mild steel from both experimental and theoretical viewpoints. Arora et al. [20] computed the strain rates and strains during the FSW of the AA2524 alloy based on a coupled viscoplastic flow and heat transfer model, and reported that the computed values of the strains rate components of materials at any locations around the welding tool and the accumulated strain components experienced by materials around the welding tool in both the welding direction and the transverse direction for an investigated FSW condition were in the ranges of 9 to 9 s1 and 10 to 5, respectively. Mironov et al. [14,16,21] have made significant efforts on clarifying the microstructural evolution of a
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variety of metallic materials during the FSW, and Liu et al. [22,23] recently proposed a new methodology to rigorously examine the microstructural evolution during FSW using an insert marker material accompanied with the “stop tool action” and liquid CO2 cooling. On the other hand, the relationship between the welding variables, microstructure and mechanical properties of the weld joints has also been extensively investigated in order to fabricate sound FSW joints by welding variable optimization. Cui et al. [24] successfully produced a high carbon steel FSW joint without any pre- or post-heat treatment by optimizing the welding variables to control the peak welding temperature and post cooling rate. Sun et al. [25] investigated the process window for the FSW of pure copper and related the final mechanical properties of copper joints with the processing dependent microstructure. However, how the initially unbonded interface of the work-pieces is joined during FSW, which is the most fundamental welding issue of FSW, still lacks full understanding. Zhou et al. [26] investigated the effect of the kissing bond, which corresponds to the initial interface of the work-pieces, on the mechanical properties and fracture behaviors of AA5083-H112 FSW joints. Depending on the welding conditions, the welds were found to tensile fracture within the stir zone along a 45 shear surface or along the kissing bond which led to a degraded mechanical property. However, the principle for the fracture along the kissing bond is still unclear. Therefore, it is apparently necessary to understand the interface joining mechanism during FSW so that the undesired fracture along the kissing bond in the FSW joints can be clearly understood and hence avoided. In this study, a butt FSW processing was performed on pure Cu and Cu-10Zn alloy plates. The reason for joining Cu-10Zn alloy to pure Cu to simulate the normal FSW is that the joining interface after FSW can be clearly identified even in microscale via element tracing in the absence of a large difference in physical and chemical properties between two materials. The microstructural evolution along the Cu/Cu-10Zn joining interface in the FSW joint was systematically investigated in detail to elucidate the interface joining mechanism during the FSW. The obtained findings not only assist us to understand the most fundamental welding principle of FSW on an atomic level, which may contribute to academic advances, but might also provide a guidance towards the joining interface microstructural control which is essential for the fabrication of sound FSW joints. 2. Experimental procedure Pure Cu and Cu-10Zn alloy plates with the dimensions of 200L 50W 3T (mm) were subjected to a butt FSW process at a tool rotation speed of 800 rpm and a welding speed of 100 mm/min under position control. The tilt angle of the tool rotation axis was set as 3 . A WC-based-alloy tool, which has a concave shoulder (12 mm diameter and 10 cone angle) and a straight cylindrical pin without any threads (4 mm diameter and 2.8 mm length) was applied in the present study. The FSW traverse was stopped prior to reaching the end of the abutment of the work-pieces and the tool was extracted. The pin-extraction-site TD (transverse direction)WD (welding direction) surface (~0.8 mm beneath the shoulder) specimen was prepared from the obtained FSW joint by an electric discharge machine as illustrated in Fig. 1. The specimen was then mechanically polished using waterproof SiC emery papers up to 4000 grits and mirror-polished using a colloidal SiO2 suspension. The mirror-polished specimen was subsequently etched using a solution of FeCl3 (5 g)þHCl (25 ml)þH2O (75 ml), and examined by an optical microscope. Fig. 2a shows the optical micrograph of the pin-extraction-site TD-WD surface specimen. It is seen that the Cu/ Cu-10Zn interface is retained and clearly visible after the FSW processing. This result shows a correspondence to the well
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Fig. 1. Schematic illustration of FSW processing of Cu/Cu-10Zn, and keyhole TD-WD surface specimen preparation for microstructural observation.
proposed material flow model of FSW [27,28] as illustrated in Fig. 2b, in which the base material ahead of the rotating tool approaches the tool, then it flows around the tool as an extrusion process of less than one rotation, and is finally left behind the tool at roughly the same transverse location. Thin foil specimens were fabricated at various positions along the visible Cu/Cu-10Zn interface by focused ion beam (FIB) as shown in Fig. 2a. Their normal directions are always lying on the plane of the TD-WD surface and parallel to the respective material flow directions as indicated by the yellow arrows in Fig. 2a. The prepared thin foil specimens were then subjected to a microstructural analysis by scanning transmission electron microscopy (STEM) and energy dispersive X-ray spectroscopy (EDS) in STEM at an accelerating voltage of 300 kV. In addition, the mechanically polished pin-extraction-site TDWD surface specimen was also electro-polished in an aqueous solution composed of 87.5% orthophosphoric acid at ~0 C (ice bath) with an applied potential of 1.5 V and then subjected to a microstructural analysis by electron backscatter diffraction (EBSD). In the present study, boundaries with misorientation angles lower than 2 were removed in the EBSD maps to eliminate spurious boundaries caused by orientation noise. The boundaries having misorientation angles ranging from 2 to 15 are defined as low-angle grain boundaries (LAGBs) shown as red lines, while the boundaries with misorientation angles 15 are denoted as high-angle grain boundaries (HAGBs) shown as black lines. 3. Results and discussion 3.1. Cu/Cu-10Zn interface microstructural evolution Fig. 3 shows the STEM-EDS results of the specimen fabricated at position 1. As illustrated in Fig. 2a, position 1 corresponds to the initial plastic deformation stage of the interface materials caused by the tool rotating and tool travelling. The Cu, O and Zn elemental mappings are shown in red, blue and green, respectively. Fig. 3e shows the magnification of the rectangle marked in Fig. 3a and the corresponding Cu, O and Zn elemental mappings are shown in Fig. 3(feh), respectively. Based on these results, it is seen that the Cu/Cu-10Zn interface is still not bonded and lots of oxides are still distributed on the interface at position 1. Fig. 4 shows the STEM-EDS results of the specimen at position 2, which corresponds to an early plastic deformation stage of the interface materials as seen in Fig. 2a. The Cu/Cu-10Zn interface can be identified based on the EDS elemental mappings and marked as
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Fig. 2. (a) Optical micrograph of the keyhole TD-WD surface specimen and thin foil specimens preparation; (b) schematic image of material flow model of FSW.
Fig. 3. STEM-EDS results of the foil specimen fabricated at position 1: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (feh) the corresponding Cu, O, Zn elemental mappings of this region, respectively.
the red dotted line in Fig. 4a and d, which indicates that the majority of the Cu/Cu-10Zn interface has adhered at this position. Small oxide particles with an average diameter of ~100 nm are dispersed on the interface based on the Zn mapping (Fig. 4d), inferring that the initial large oxides were fragmented into small particles and then distributed along the interface at this position. An interface region marked by the yellow rectangle in Fig. 4a is magnified in Fig. 4e. It is noted that a grain boundary bulging occurs between two small oxide particles, as indicated by the yellow
dotted line in Fig. 4e. The region swept by the migrating boundary is seen to lack dislocations. Furthermore, a high density of tangled dislocations is observed in the convex side of the migrating boundary within the Cu-10Zn as indicated by the black arrows in Fig. 4a and e, whereas almost no dislocations can be identified on the concave side of the migrating boundary within the Cu. Similar phenomenon can be observed in the EBSD results derived near position 2 as shown in Fig. 5. The Cu/Cu-10Zn interface can be identified based on the locations of the oxide particles as indicated
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Fig. 4. STEM-EDS results of the foil specimen derived at position 2: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (feh) the corresponding Cu, O, Zn elemental mappings of this region, respectively.
Fig. 5. (aeb) EBSD results derived from the location near position 2; (c) schematic model for strain-induced grain boundary migration.
by the red arrows in the image quality (IQ) map and marked as the yellow dotted line in the inverse pole figure (IPF) map. The interface regions A and B circled by the solid and dotted lines, respectively, in Fig. 5a are found to show a similar oxide particle distribution, but different boundary behaviors; the boundary in region A shows an apparent bulging towards the Cu-10Zn side, whereas the boundary in region B does not obviously bulge. Furthermore, in region A, LAGBs composed of well rearranged dislocations are observed on the convex side of the bulged boundary within the Cu-10Zn while no LAGB exists on the concave side within the Cu. In contrast, no LAGB can be observed on both sides of the boundary in region B. These results show a consistency with those obtained from the above STEM-EDS data (Fig. 4). According to all the above results, this grain boundary bulging is thus identified as the strain-induced grain boundary migration [29], which is driven by the stored energy difference associated with the difference in the dislocation contents between the opposite sides of the grain boundary. On the other hand, it is also worth noting that a void is visible adjacent to
oxide particle 1, whereas no void is detectable besides oxide particle 2, based on the STEM-EDS data shown in Fig. 4eeh. Fig. 6 shows the STEM-EDS results of the specimen at position 3. The initial Cu/Cu-10Zn interface is identified based on the elemental mappings and marked as the red dotted lines in Fig. 6a and d. Small oxide particles (~100 nm in diameter) as well as a relatively large void (~250 nm in diameter) are observed on the interface, as indicated by the white arrows in Fig. 6d. The formation of the large void is considered to be attributed to the drop out of the large oxide particle, which initially existed at this location, during the FIB machining. Fig. 6e shows the magnification of the rectangle marked in Fig. 6a. Selected area electron diffraction (SAED) patterns of locations 1-6 marked in Fig. 6e were derived with a constant beam direction and are exhibited in Fig. 6(1e6), respectively. Locations 1e3 show the same SAED patterns, suggesting that they are contained in one individual grain and so are locations 4 and 5. The grain boundaries are thus clearly identified and emphasized as the yellow dotted lines in Fig. 6e. These results indicate that the strain-
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Fig. 6. STEM-EDS results of the foil specimen derived at position 3: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (1e6) selected area electron diffraction patterns derived from areas marked by circles 1e6 in (e), respectively.
induced grain boundary migration, which is pinned by oxide particles, has developed at this position under the further deformation. Fig. 7 shows the STEM-EDS results of the specimen located at position 4. The initial Cu/Cu-10Zn interface is identified as the red dotted lines drawn in Fig. 7a and d based on the elemental mappings. The two areas located on opposite sides of the Cu/Cu-10Zn interface marked by the circle A and the circle B in Fig. 7a show the same SAED pattern, i.e., the same crystal orientation, inferring that a much more remarkable strain-induced grain boundary migration has occurred at this position as illustrated by the yellow dotted line in Fig. 7a. Fig. 7e shows the magnification of the green rectangle marked in Fig. 7a. The Zn boundary, which is originally considered as the Cu/Cu-10Zn interface, is marked by the red dotted lines in Fig. 7e and h. A STEM-EDS line analysis was conducted along the red arrow marked in Fig. 7e and the results are shown in Figs. 7(e1). It is noted that the Zn element has a longrange diffusion zone of around 0.4 mm which is bounded by the green dotted lines in Figs. 7(e1), and thus the real initial Cu/Cu10Zn interface is likely to be located in the central region of the diffusion zone as marked by the red dotted line in Figs. 7(e1), which corresponds to the location marked by the red solid line on the arrow in Fig. 7e. In addition, small oxide particles and voids are observed based on the elemental mappings and their locations in principle indicate the position of the initial Cu/Cu-10Zn interface. Based on these results, the real initial Cu/Cu-10Zn interface can be reasonably identified as the green dotted line in Fig. 7e. Moreover, it is worth noting that an individual grain, which crosses the Cu/Cu10Zn interface and is free of dislocation, can be detected in Fig. 7e as illustrated by the yellow dotted line. Lots of small oxide particles are dispersed inside the grain. It is considered that in the region rich of small oxide particles, dislocations can be strongly pinned by these oxide particles and accumulate and tangle around them, thus
a highly-strained deformation zone is produced. A nucleus free of dislocations thus tends to form inside this highly-strained deformation zone, and grows via HAGB migration along with consuming the surrounding highly-strained matrix. Fig. 7i shows magnification of the yellow rectangle marked in Fig. 7a. Similar to the analysis in Fig. 7e, the Zn boundary can be identified by the red dotted lines in Fig. 7i and l based on the elemental mappings, while the real initial Cu/Cu-10Zn interface can be identified by the green dotted line in Fig. 7i according to the EDS line analysis (Figs. 7(i-1)) associated with the Zn element diffusion and the oxide particle positions. The green dotted line, i.e., the initial Cu/Cu-10Zn interface, is observed to overlap with some segments of the grain boundaries, which suggests that these segments of grain boundaries are the initial Cu/Cu-10Zn interface, and have not yet migrated or cultivated any new grains up to position 4. In addition, the large voids correlated with the large oxide particles are still retained on the interface at this position. Based on these results, it is known that some segments of the initial Cu/Cu-10Zn interface have migrated, driven by the stored energy difference between opposite sides of the interface, whereas some others have not yet migrated. A HRTEM analysis was thus conducted on the non-migrated and migrated interface segments in order to comparatively investigate the atomic bonding conditions for both types of interfaces. Fig. 8 shows the HRTEM images of the non-migrated (Fig. 8b) and migrated (Fig. 8d) interface segments, derived from the regions circled in red marked in Fig. 8a and c, respectively. Fig. 8c is the magnified image of the region marked by the black rectangle shown in Fig. 7e which contains the migrated interface segments indicated by the black dotted curve. The observed non-migrated Cu/Cu-10Zn interface shows a semicoherent interface structure and no oxide crystallographic structure can be identified on the interface (Fig. 8b), which indicates that
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Fig. 7. STEM-EDS results of the foil specimen derived at position 4: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by green dotted rectangle shown in (a); (feh) corresponding Cu, O, Zn elemental mappings, respectively; (i) magnification of the region marked by yellow dotted rectangle shown in (a); (jel) corresponding Cu, O, Zn elemental mappings, respectively; (e1 and i-1) elemental distribution along the red arrows marked in (e) and (i), respectively, based on EDS line analysis. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
fresh surfaces were produced and atomic-level joining was achieved at this interface. For the observed migrated Cu/Cu-10Zn interface, it is also identified that no oxide crystallographic information is detected and the atomic bonding has been achieved at the interface (Fig. 8d). This result suggests that the fresh surfaces
seemed to be first produced and atomically joined to form a grain boundary at this interface; then, the driving force for the grain boundary migration was achieved by different dislocation accumulation rates between the opposite sides of the boundary and the grain boundary migration, i.e., interface migration, occurred at this
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Fig. 8. (a) Fig. 7i which contains the non-migrated interface segments; (b) HRTEM image of the non-migrated interface segment region circled by red in (a); (c) magnification of the region marked by black rectangle shown in Fig. 7e; (d) HRTEM image of the migrated interface segment region circled by red in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
interface. Fig. 9 shows the STEM-EDS results of the specimen derived at position 5. It is found that no large oxide particle is visible on the Cu/Cu-10Zn interface and all the grain boundaries initially on the Cu/Cu-10Zn interface have migrated at this position. Fig. 9e shows the magnification of the rectangle drawn in Fig. 9a. It is noted that a small oxide particle with a diameter of approximately 60 nm is sporadically visible in the grain interior. 3.2. Discussion The results obtained from Fig. 3 revealed that the initial Cu/Cu10Zn interface is completely unbonded and lots of oxides are densely distributed at the interface at position 1, i.e., the initial plastic deformation stage of the interface materials around the tool during the FSW. It is not surprising that the oxide regions in Fig. 3 show a higher O contrast in the O mappings and the comparatively lower Cu contrast in the Cu mappings because even the Cu-10Zn region has a high Cu content of ~90 at%. However, in the Zn mappings, the oxide regions show a much lower Zn contrast compared to the Cu-10Zn region. These results indicate that the oxides initially distributed on the Cu/Cu-10Zn interface at position 1 are mainly Cu oxides. However, when the interface evolved to position 2, i.e., an early plastic deformation stage of the interface materials, the small oxide particles distributed along the interface are seen to show the highest Zn contrast in the Zn mappings. All the oxide particles observed on the interface at the positions following position 2 are found to show similar characteristics. In Fig. 7i at position 4, an oxide particle detected on the interface marked by the yellow dotted circle is identified to be composed of ~29.1 at% O, ~42.8 at% Cu, and ~28.1 at% Zn based on the STEM-EDS spot analysis. Therefore, it is concluded that the oxide particles observed at
the interface positions after position 2 are likely a mixture of Cu oxides and Zn oxides. It is well known that the Cu oxides can react with Zn to form Zn oxides and Cu at high temperature because of the different reactivities of Cu and Zn, accompanied by a release of heat. Thus, it is inferred that during the interface material flow deformation around the tool in the FSW, the Cu oxides, which initially formed on the interface, reacted with Zn at an early plastic deformation stage, then a mixed oxide consisting of Cu oxide and Zn oxide finally formed on the interface with a reaction heat release. These findings might provide a new insight into taking advantage of the chemical reaction associated with an endothermic or exothermic behavior that occurred on the joining interface to possibly control the microstructure of the joining interface. The occurrence of the strain-induced grain boundary migration, which was pinned by the oxide particles, has been frequently identified on the Cu/Cu-10Zn joining interface during the FSW. The driving force for the boundary migration is regarded as the stored energy difference correlated to the dislocation density difference between the opposite sides of the boundary. Fig. 5c shows a schematic model for the strain-induced grain boundary migration [29]. Grain A and grain B have different dislocation densities, rA and rB , respectively. The DE, which is the stored energy difference between the opposite sides of the grain boundary, can thus be calculated according to the following equation: DE ¼ 12 GðrB rA Þb2 , where G is the shear modulus of the material, and b is the burgers vector of a dislocation. The bulging boundary shows a spherical cap shape with a radius of R and has a specific grain boundary surface energy of d. The condition for the grain boundary migration to occur is the favorable energy balance between the decrease in the total stored energy due to dislocation elimination and rearrangement caused by the passing boundary and the increase in total grain boundary surface energy due to the newly bulged boundary. This occurrence
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Fig. 9. STEM-EDS results of the foil specimen derived at position 5: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (feh) corresponding Cu, O, Zn elemental mappings of this region, respectively.
condition is thus given by the following formula: DE > 2Ld, where L is half of the distance between the intersections of the bulging boundary and the parent boundary. In the present study, since the strain-induced grain boundary migration is generally pinned by the distributed oxide particles, L can be regarded as half of the distance between the neighboring oxide particles. Therefore, it is concluded that the occurrence of the strain-induced grain boundary migration on the Cu/Cu-10Zn joining interface in the present study can be influenced by the following three factors, the stored energy difference between the opposite sides of the grain boundary, the specific grain boundary surface energy, and the interspacing between the oxide particles distributed on the interface. This view is well supported by the EBSD results shown in Fig. 10 which exhibits the Cu/Cu-10Zn joining interface microstructure derived from the location near position 3. The joining interface can be identified according to the distribution of the oxide particles indicated by the red arrows in the IQ map (Fig. 10b) and marked as the red dotted line in the IPF map (Fig. 10a). The interface regions indicated by the rectangles A, B, and C have different morphologies and distribution of the oxide particles thereby showing different interface behaviors. In region A, since almost no oxide particle is visible on the interface, the grain boundary migration occurs at this location; while for region B, no grain boundary migration is identified because lots of relatively large oxide particles are densely distributed on the interface leading to a small interspacing between the oxide particles. It is worth noting that even though the oxide particles are observed to be dispersed on the interface in region C, since the size and number density of the oxide particles are relatively small, the grain boundary migration can still occur at this position. In addition, it has been noted that the void is visible neighboring oxide particle 1, whereas no void is detectable beside oxide particle 2 according to Fig. 4eeh. Furthermore, it can also be seen that the strain-induced grain boundary migration occurring between these two oxide particles has overcome the pinning of oxide particle 2 but
still been pinned by oxide particle 1. These results suggest that the voids could form on the joining interface neighboring the oxide particles during the FSW because the existence of the oxide particles impedes the complete and ideal adhesion of the interface. The strain-induced grain boundary migration can promote the atomic diffusion and dislocation movement, rearrangement and annihilation in the nearby of the boundary, which will help to extinguish the retained voids adjacent to the oxide particles. Therefore, when the strain-induced grain boundary migration overcomes the pinning effect and passes through the entire oxide particles, the voids retained in the surroundings of the oxide particles are likely to be extinguished completely. It is also worth noting that the strain-induced grain boundary migration observed in the present study occurs from the Cu side towards the Cu-10Zn side. The reason for this phenomenon can be demonstrated as follows. Cu-10Zn is known to have a lower stacking fault energy of ~35 mJ/m2 compared to that of pure Cu (~78 mJ/m2) because of the alloying of Zn element [30]. Therefore, the full dislocations tend to dissociate into the partial dislocations and the dislocation cross-slip and dislocation climb become more difficult on the Cu-10Zn side, which makes the dislocation recovery associated with the dislocation mobility, rearrangement and annihilation more difficult thereby promoting the dislocation accumulation in the Cu-10Zn side compared with the Cu side. Therefore, it is easy for the Cu-10Zn side to exhibit a higher dislocation density compared to the Cu side, which causes the strain-induced grain boundary migration to easily occur from the Cu side towards the Cu-10Zn side due to the favorable stored energy balance. In the present study, the stacking fault energy difference between two work-pieces plays an important role on producing the stored energy difference between the opposite sides of the joining interface. However, it should be noted that even under the lack of stacking fault energy difference, the stored energy difference will also generate because the different grain orientations and different stress conditions between the opposite sides of the joining
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Fig. 10. EBSD results derived from the location near position 3.
interface can provide different dislocation generation rates and accumulation rates. In addition, it is seen that the apparent diffusion of Zn element across the Cu/Cu-10Zn interface can be identified in both the locations indicated by the red arrows in Fig. 7e and i based on the EDS line analysis of Fig. 7ee1 and i-1, and the location in Fig. 7e shows a relatively longer Zn diffusion distance of ~0.4 mm compared to that of the location in Fig. 7i (~0.25 mm). This difference in the Zn diffusion distance is considered to be associated with the fresh surface and atomic-level interface formation caused by the intense shear and compression as well as the thermal cycles and dwells at a sufficient temperature at these two locations during the FSW. In addition, it has been mentioned that the dislocations were strongly pinned and accumulated to form a highly-strained deformation zone at the interface in the location of Fig. 7e that facilitated the formation and growth of the dislocation-free nucleus inside it, while the interface in the location of Fig. 7i has not yet migrated. Both the formation of the highly-strained zone consisting of abundant defects and the formation and growth of the new grain accompanied with dislocation rearrangement and annihilation can promote the diffusion of the Zn element across the interface in the location of Fig. 7e, which can be one of the possible explanations for the relatively longer Zn diffusion distance in the location of Fig. 7e. Based on these results, the interface joining mechanism during the FSW can be concluded as illustrated in Fig. 11. For the initial joining interface of the work-pieces, it is completely unbonded and lots of large oxides are distributed along the initial joining interface as shown in Fig. 11a. The large oxides are then fragmented into small particles, and the oxide layers formed on the interface are ruptured due to the compressive force perpendicular to the interface and the shear force and interface material deformation parallel to the interface, which produces fresh surfaces on both Cu side and Cu-10Zn side as illustrated in Fig. 11b. Then an atomic-level joining is achieved between the fresh surfaces of both Cu and Cu-10Zn sides under the compressive force, while leaving the small oxide particles still dispersed on the interface with the voids formed around the oxide particles. As demonstrated in the magnified schematic inset in Fig. 11c, the dislocation slip and/or atomic diffusion occurring around the void assists to extinguish the void.
However, it is believed that a large deformation and/or high temperature is required to provide extensive dislocation gliding and/or abundant atomic diffusion around the oxide particles in order to completely extinguish the retained voids neighboring the oxide particles. For this interface condition, the stress concentration easily occurs at the locations of the oxide particles which may initiate the micro-cracks. The micro-cracks then easily propagate along the interface and connect to each other, then finally leading to a fracture along the interface. In fact, a strain-induced grain boundary migration will occur on the interface, which is driven by the stored energy difference associated with the dislocation density difference between the opposite sides of the grain boundary. This grain boundary migration will then promote the dislocation movement and atomic diffusion near the migrating boundary thus assisting to extinguish the voids formed around the oxide particles and also shift the locations of the oxide particles from the grain boundaries to grain interiors, as illustrated in Fig. 11d and e. Under further deformation, the oxide particles on the interface will be further fragmented into fine particles and further dispersed; finally, a sound interface joining without oxide particles and large voids remaining on the interface can be obtained during the FSW as demonstrated in Fig. 11f. 4. Summary In this study, a butt FSW process was performed using Cu and Cu-10Zn plates. Microstructural evolution along the Cu/Cu-10Zn interface in the fabricated FSW joint was systematically investigated to clarify the interface joining mechanism during the FSW. The obtained findings are shown as follows: (1) During the FSW, the large oxides initially existing on the work-pieces joining interface are fragmented into small oxide particles and the oxide layer is ruptured under the compressive force perpendicular to the interface and the shear force and interface material deformation parallel to the interface, which produces fresh surfaces on the interface. (2) The newly formed fresh surfaces of the work-pieces are then joined on an atomic-level under the compressive force, while
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Fig. 11. Schematic illustration of interface joining mechanism during FSW.
the small oxide particles are still dispersed on the interface with the voids formed around the particles. (3) A strain-induced grain boundary migration driven by the stored energy difference between the opposite sides of the boundary will actually occur at the interface, which facilitates extinguishing the voids formed around the oxide particles and shifts the locations of the oxide particles from the interface consisting of grain boundaries to the grain interiors. (4) The oxide particles are further fragmented and dispersed, and the strain-induced grain boundary migration is further developed under deformation, hence a sound interface joining without oxide particles and voids remaining on the interface is finally obtained during the FSW. Acknowledgements The authors wish to acknowledge the financial support by the New Energy and Industrial Technology Development Organization (NEDO) under the “Innovation Structural Materials Project (Future Pioneering Projects)”, JSPS KAKENHI Grant Numbers JP15H04133 and JP18K14027, and a ISIJ Research Promotion Grant. References [1] W.M. Thomas, E.D. Nicholas, J.C. Needham, M.G. Murch, P. Templesmith, C.J. Dawes, G.B. Patent Application No. 9125978.8 (1991). [2] R.S. Mishra, Z.Y. Ma, Friction stir welding and processing, Mater. Sci. Eng. R 50 (2005) 1e78. [3] C.G. Rhodes, M.W. Mahoney, W.H. Bingel, M. Calabrese, Fine-grain evolution in friction-stir processed 7050 aluminum, Scripta Mater. 48 (2003) 1451e1455. [4] K. Kitamura, H. Fujii, Y. Iwata, Y.S. Sun, Y. Morisada, Flexible control of the microstructure and mechanical properties of friction stir welded Ti-6Al-4V joints, Mater. Des. 46 (2013) 348e354. [5] S. Mironov, K. Inagaki, Y.S. Sato, H. Kokawa, Effect of welding temperature on
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[21] S. Mironov, T. Onuma, Y.S. Sato, H. Kokawa, Microstructure evolution during friction-stir welding of AZ31 magnesium alloy, Acta Mater. 100 (2015) 301e312. [22] X.C. Liu, Y.F. Sun, T. Nagira, H. Fujii, Investigation of temperature dependent microstructure evolution of pure iron during friction stir welding using liquid CO2 rapid cooling, Mater. Char. 137 (2018) 24e38. [23] X.C. Liu, Y.F. Sun, T. Nagira, K. Ushioda, H. Fujii, Microstructure evolution of Cu-30Zn during friction stir welding, J. Mater. Sci. 53 (2018) 10423e10441. [24] L. Cui, H. Fujii, N. Tsuji, K. Nogi, Friction stir welding of a high carbon steel, Scripta Mater. 56 (2007) 637e640. [25] Y.F. Sun, H. Fujii, Investigation of the welding parameter dependent microstructure and mechanical properties of friction stir welded pure copper, Mater. Sci. Eng. 527 (2010) 6879e6886.
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Full length article
Elucidation of interface joining mechanism during friction stir welding through Cu/Cu-10Zn interfacial observations Huihong Liu*, Kohsaku Ushioda, Hidetoshi Fujii Joining and Welding Research Institute, Osaka University, Ibaraki, Osaka, 567-0047, Japan
a r t i c l e i n f o
a b s t r a c t
Article history: Received 1 October 2018 Received in revised form 31 December 2018 Accepted 3 January 2019 Available online 4 January 2019
A butt friction stir welding (FSW) process was performed on Cu and Cu-10Zn plates. Microstructural evolution along the Cu/Cu-10Zn interface retained in the FSW joint was systematically investigated in order to clarify the interface joining mechanism during the FSW. The initial joining interface is completely unbonded and has lots of large oxides distributed on it. The initial large oxides are then fragmented into small particles and the oxide layers initially formed on the interface are ruptured, which produces fresh surfaces on both the Cu and Cu-10Zn sides. The fresh surfaces are then joined at an atomic level under a compressive force perpendicular to the interface, while the small oxide particles are still dispersed on the interface with voids formed around the particles. A strain-induced grain boundary migration driven by the stored energy difference associated with the dislocation density difference will occur, which assists to extinguish the voids retained beside the oxide particles. Finally, a sound interface joining without oxide particles and large voids remaining on the interface is obtained during the FSW. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Friction stir welding STEM Microstructure Grain boundary migration Interface joining mechanism
1. Introduction Friction stir welding (FSW) is an innovative joining technique, which was invented by The Welding Institute (TWI) of the UK in 1991 originally aiming to join Al and its alloys that are difficult to fusion weld [1]. In FSW, a non-consumable rotating tool with the designated shoulder and pin is inserted into the interface of two butting work-pieces. The heat is mainly generated by the friction between the tool and the work-pieces, and it softens the workpiece materials around the tool, making them plastically deformable. The rotating tool then traverses along the butting line of the work-pieces accompanied by transporting the surrounding materials from the leading side to the trailing side of the tool to produce a joint [2]. Since FSW enables a joining in the solid state, it can be adopted to avoid or minimize the fusion-welding-associated problems such as formation of coarse columnar grains, high residual stress, severe welding distortion, and solidification defects as well as macro-segregation, which inevitably degrade the quality of the fusion weld joints. Moreover, since a fine and equiaxed recrystallized microstructure can be obtained in the friction stir welds by severe plastic deformation and elevated temperature, and
* Corresponding author. E-mail address: [email protected] (H. Liu). https://doi.org/10.1016/j.actamat.2019.01.004 1359-6454/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
the formation of a heat-affected-zone, which is generally the fracture site in a fusion-welded joint, can be effectively suppressed by the relatively low welding temperature, sound weld joints with good mechanical properties can be produced by the FSW [3e5]. Nowadays, FSW, which is considered as a green and nextgeneration joining technology, has been extensively applied in low-melting-point materials, such as Al and Mg alloys [6e10], and has also been expanded to many other high-melting-point materials including Cu, Fe and Ti alloys due to the development of highly durable tools [11e18]. Most of the FSW researches have focused on clarifying the physical fundamentals of FSW such as heat transfer, material flow, strain and strain rate, and microstructural evolution during FSW. Nandan et al. [19] investigated the three-dimensional heat transfer and material flow during FSW of mild steel from both experimental and theoretical viewpoints. Arora et al. [20] computed the strain rates and strains during the FSW of the AA2524 alloy based on a coupled viscoplastic flow and heat transfer model, and reported that the computed values of the strains rate components of materials at any locations around the welding tool and the accumulated strain components experienced by materials around the welding tool in both the welding direction and the transverse direction for an investigated FSW condition were in the ranges of 9 to 9 s1 and 10 to 5, respectively. Mironov et al. [14,16,21] have made significant efforts on clarifying the microstructural evolution of a
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variety of metallic materials during the FSW, and Liu et al. [22,23] recently proposed a new methodology to rigorously examine the microstructural evolution during FSW using an insert marker material accompanied with the “stop tool action” and liquid CO2 cooling. On the other hand, the relationship between the welding variables, microstructure and mechanical properties of the weld joints has also been extensively investigated in order to fabricate sound FSW joints by welding variable optimization. Cui et al. [24] successfully produced a high carbon steel FSW joint without any pre- or post-heat treatment by optimizing the welding variables to control the peak welding temperature and post cooling rate. Sun et al. [25] investigated the process window for the FSW of pure copper and related the final mechanical properties of copper joints with the processing dependent microstructure. However, how the initially unbonded interface of the work-pieces is joined during FSW, which is the most fundamental welding issue of FSW, still lacks full understanding. Zhou et al. [26] investigated the effect of the kissing bond, which corresponds to the initial interface of the work-pieces, on the mechanical properties and fracture behaviors of AA5083-H112 FSW joints. Depending on the welding conditions, the welds were found to tensile fracture within the stir zone along a 45 shear surface or along the kissing bond which led to a degraded mechanical property. However, the principle for the fracture along the kissing bond is still unclear. Therefore, it is apparently necessary to understand the interface joining mechanism during FSW so that the undesired fracture along the kissing bond in the FSW joints can be clearly understood and hence avoided. In this study, a butt FSW processing was performed on pure Cu and Cu-10Zn alloy plates. The reason for joining Cu-10Zn alloy to pure Cu to simulate the normal FSW is that the joining interface after FSW can be clearly identified even in microscale via element tracing in the absence of a large difference in physical and chemical properties between two materials. The microstructural evolution along the Cu/Cu-10Zn joining interface in the FSW joint was systematically investigated in detail to elucidate the interface joining mechanism during the FSW. The obtained findings not only assist us to understand the most fundamental welding principle of FSW on an atomic level, which may contribute to academic advances, but might also provide a guidance towards the joining interface microstructural control which is essential for the fabrication of sound FSW joints. 2. Experimental procedure Pure Cu and Cu-10Zn alloy plates with the dimensions of 200L 50W 3T (mm) were subjected to a butt FSW process at a tool rotation speed of 800 rpm and a welding speed of 100 mm/min under position control. The tilt angle of the tool rotation axis was set as 3 . A WC-based-alloy tool, which has a concave shoulder (12 mm diameter and 10 cone angle) and a straight cylindrical pin without any threads (4 mm diameter and 2.8 mm length) was applied in the present study. The FSW traverse was stopped prior to reaching the end of the abutment of the work-pieces and the tool was extracted. The pin-extraction-site TD (transverse direction)WD (welding direction) surface (~0.8 mm beneath the shoulder) specimen was prepared from the obtained FSW joint by an electric discharge machine as illustrated in Fig. 1. The specimen was then mechanically polished using waterproof SiC emery papers up to 4000 grits and mirror-polished using a colloidal SiO2 suspension. The mirror-polished specimen was subsequently etched using a solution of FeCl3 (5 g)þHCl (25 ml)þH2O (75 ml), and examined by an optical microscope. Fig. 2a shows the optical micrograph of the pin-extraction-site TD-WD surface specimen. It is seen that the Cu/ Cu-10Zn interface is retained and clearly visible after the FSW processing. This result shows a correspondence to the well
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Fig. 1. Schematic illustration of FSW processing of Cu/Cu-10Zn, and keyhole TD-WD surface specimen preparation for microstructural observation.
proposed material flow model of FSW [27,28] as illustrated in Fig. 2b, in which the base material ahead of the rotating tool approaches the tool, then it flows around the tool as an extrusion process of less than one rotation, and is finally left behind the tool at roughly the same transverse location. Thin foil specimens were fabricated at various positions along the visible Cu/Cu-10Zn interface by focused ion beam (FIB) as shown in Fig. 2a. Their normal directions are always lying on the plane of the TD-WD surface and parallel to the respective material flow directions as indicated by the yellow arrows in Fig. 2a. The prepared thin foil specimens were then subjected to a microstructural analysis by scanning transmission electron microscopy (STEM) and energy dispersive X-ray spectroscopy (EDS) in STEM at an accelerating voltage of 300 kV. In addition, the mechanically polished pin-extraction-site TDWD surface specimen was also electro-polished in an aqueous solution composed of 87.5% orthophosphoric acid at ~0 C (ice bath) with an applied potential of 1.5 V and then subjected to a microstructural analysis by electron backscatter diffraction (EBSD). In the present study, boundaries with misorientation angles lower than 2 were removed in the EBSD maps to eliminate spurious boundaries caused by orientation noise. The boundaries having misorientation angles ranging from 2 to 15 are defined as low-angle grain boundaries (LAGBs) shown as red lines, while the boundaries with misorientation angles 15 are denoted as high-angle grain boundaries (HAGBs) shown as black lines. 3. Results and discussion 3.1. Cu/Cu-10Zn interface microstructural evolution Fig. 3 shows the STEM-EDS results of the specimen fabricated at position 1. As illustrated in Fig. 2a, position 1 corresponds to the initial plastic deformation stage of the interface materials caused by the tool rotating and tool travelling. The Cu, O and Zn elemental mappings are shown in red, blue and green, respectively. Fig. 3e shows the magnification of the rectangle marked in Fig. 3a and the corresponding Cu, O and Zn elemental mappings are shown in Fig. 3(feh), respectively. Based on these results, it is seen that the Cu/Cu-10Zn interface is still not bonded and lots of oxides are still distributed on the interface at position 1. Fig. 4 shows the STEM-EDS results of the specimen at position 2, which corresponds to an early plastic deformation stage of the interface materials as seen in Fig. 2a. The Cu/Cu-10Zn interface can be identified based on the EDS elemental mappings and marked as
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Fig. 2. (a) Optical micrograph of the keyhole TD-WD surface specimen and thin foil specimens preparation; (b) schematic image of material flow model of FSW.
Fig. 3. STEM-EDS results of the foil specimen fabricated at position 1: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (feh) the corresponding Cu, O, Zn elemental mappings of this region, respectively.
the red dotted line in Fig. 4a and d, which indicates that the majority of the Cu/Cu-10Zn interface has adhered at this position. Small oxide particles with an average diameter of ~100 nm are dispersed on the interface based on the Zn mapping (Fig. 4d), inferring that the initial large oxides were fragmented into small particles and then distributed along the interface at this position. An interface region marked by the yellow rectangle in Fig. 4a is magnified in Fig. 4e. It is noted that a grain boundary bulging occurs between two small oxide particles, as indicated by the yellow
dotted line in Fig. 4e. The region swept by the migrating boundary is seen to lack dislocations. Furthermore, a high density of tangled dislocations is observed in the convex side of the migrating boundary within the Cu-10Zn as indicated by the black arrows in Fig. 4a and e, whereas almost no dislocations can be identified on the concave side of the migrating boundary within the Cu. Similar phenomenon can be observed in the EBSD results derived near position 2 as shown in Fig. 5. The Cu/Cu-10Zn interface can be identified based on the locations of the oxide particles as indicated
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Fig. 4. STEM-EDS results of the foil specimen derived at position 2: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (feh) the corresponding Cu, O, Zn elemental mappings of this region, respectively.
Fig. 5. (aeb) EBSD results derived from the location near position 2; (c) schematic model for strain-induced grain boundary migration.
by the red arrows in the image quality (IQ) map and marked as the yellow dotted line in the inverse pole figure (IPF) map. The interface regions A and B circled by the solid and dotted lines, respectively, in Fig. 5a are found to show a similar oxide particle distribution, but different boundary behaviors; the boundary in region A shows an apparent bulging towards the Cu-10Zn side, whereas the boundary in region B does not obviously bulge. Furthermore, in region A, LAGBs composed of well rearranged dislocations are observed on the convex side of the bulged boundary within the Cu-10Zn while no LAGB exists on the concave side within the Cu. In contrast, no LAGB can be observed on both sides of the boundary in region B. These results show a consistency with those obtained from the above STEM-EDS data (Fig. 4). According to all the above results, this grain boundary bulging is thus identified as the strain-induced grain boundary migration [29], which is driven by the stored energy difference associated with the difference in the dislocation contents between the opposite sides of the grain boundary. On the other hand, it is also worth noting that a void is visible adjacent to
oxide particle 1, whereas no void is detectable besides oxide particle 2, based on the STEM-EDS data shown in Fig. 4eeh. Fig. 6 shows the STEM-EDS results of the specimen at position 3. The initial Cu/Cu-10Zn interface is identified based on the elemental mappings and marked as the red dotted lines in Fig. 6a and d. Small oxide particles (~100 nm in diameter) as well as a relatively large void (~250 nm in diameter) are observed on the interface, as indicated by the white arrows in Fig. 6d. The formation of the large void is considered to be attributed to the drop out of the large oxide particle, which initially existed at this location, during the FIB machining. Fig. 6e shows the magnification of the rectangle marked in Fig. 6a. Selected area electron diffraction (SAED) patterns of locations 1-6 marked in Fig. 6e were derived with a constant beam direction and are exhibited in Fig. 6(1e6), respectively. Locations 1e3 show the same SAED patterns, suggesting that they are contained in one individual grain and so are locations 4 and 5. The grain boundaries are thus clearly identified and emphasized as the yellow dotted lines in Fig. 6e. These results indicate that the strain-
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Fig. 6. STEM-EDS results of the foil specimen derived at position 3: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (1e6) selected area electron diffraction patterns derived from areas marked by circles 1e6 in (e), respectively.
induced grain boundary migration, which is pinned by oxide particles, has developed at this position under the further deformation. Fig. 7 shows the STEM-EDS results of the specimen located at position 4. The initial Cu/Cu-10Zn interface is identified as the red dotted lines drawn in Fig. 7a and d based on the elemental mappings. The two areas located on opposite sides of the Cu/Cu-10Zn interface marked by the circle A and the circle B in Fig. 7a show the same SAED pattern, i.e., the same crystal orientation, inferring that a much more remarkable strain-induced grain boundary migration has occurred at this position as illustrated by the yellow dotted line in Fig. 7a. Fig. 7e shows the magnification of the green rectangle marked in Fig. 7a. The Zn boundary, which is originally considered as the Cu/Cu-10Zn interface, is marked by the red dotted lines in Fig. 7e and h. A STEM-EDS line analysis was conducted along the red arrow marked in Fig. 7e and the results are shown in Figs. 7(e1). It is noted that the Zn element has a longrange diffusion zone of around 0.4 mm which is bounded by the green dotted lines in Figs. 7(e1), and thus the real initial Cu/Cu10Zn interface is likely to be located in the central region of the diffusion zone as marked by the red dotted line in Figs. 7(e1), which corresponds to the location marked by the red solid line on the arrow in Fig. 7e. In addition, small oxide particles and voids are observed based on the elemental mappings and their locations in principle indicate the position of the initial Cu/Cu-10Zn interface. Based on these results, the real initial Cu/Cu-10Zn interface can be reasonably identified as the green dotted line in Fig. 7e. Moreover, it is worth noting that an individual grain, which crosses the Cu/Cu10Zn interface and is free of dislocation, can be detected in Fig. 7e as illustrated by the yellow dotted line. Lots of small oxide particles are dispersed inside the grain. It is considered that in the region rich of small oxide particles, dislocations can be strongly pinned by these oxide particles and accumulate and tangle around them, thus
a highly-strained deformation zone is produced. A nucleus free of dislocations thus tends to form inside this highly-strained deformation zone, and grows via HAGB migration along with consuming the surrounding highly-strained matrix. Fig. 7i shows magnification of the yellow rectangle marked in Fig. 7a. Similar to the analysis in Fig. 7e, the Zn boundary can be identified by the red dotted lines in Fig. 7i and l based on the elemental mappings, while the real initial Cu/Cu-10Zn interface can be identified by the green dotted line in Fig. 7i according to the EDS line analysis (Figs. 7(i-1)) associated with the Zn element diffusion and the oxide particle positions. The green dotted line, i.e., the initial Cu/Cu-10Zn interface, is observed to overlap with some segments of the grain boundaries, which suggests that these segments of grain boundaries are the initial Cu/Cu-10Zn interface, and have not yet migrated or cultivated any new grains up to position 4. In addition, the large voids correlated with the large oxide particles are still retained on the interface at this position. Based on these results, it is known that some segments of the initial Cu/Cu-10Zn interface have migrated, driven by the stored energy difference between opposite sides of the interface, whereas some others have not yet migrated. A HRTEM analysis was thus conducted on the non-migrated and migrated interface segments in order to comparatively investigate the atomic bonding conditions for both types of interfaces. Fig. 8 shows the HRTEM images of the non-migrated (Fig. 8b) and migrated (Fig. 8d) interface segments, derived from the regions circled in red marked in Fig. 8a and c, respectively. Fig. 8c is the magnified image of the region marked by the black rectangle shown in Fig. 7e which contains the migrated interface segments indicated by the black dotted curve. The observed non-migrated Cu/Cu-10Zn interface shows a semicoherent interface structure and no oxide crystallographic structure can be identified on the interface (Fig. 8b), which indicates that
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Fig. 7. STEM-EDS results of the foil specimen derived at position 4: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by green dotted rectangle shown in (a); (feh) corresponding Cu, O, Zn elemental mappings, respectively; (i) magnification of the region marked by yellow dotted rectangle shown in (a); (jel) corresponding Cu, O, Zn elemental mappings, respectively; (e1 and i-1) elemental distribution along the red arrows marked in (e) and (i), respectively, based on EDS line analysis. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
fresh surfaces were produced and atomic-level joining was achieved at this interface. For the observed migrated Cu/Cu-10Zn interface, it is also identified that no oxide crystallographic information is detected and the atomic bonding has been achieved at the interface (Fig. 8d). This result suggests that the fresh surfaces
seemed to be first produced and atomically joined to form a grain boundary at this interface; then, the driving force for the grain boundary migration was achieved by different dislocation accumulation rates between the opposite sides of the boundary and the grain boundary migration, i.e., interface migration, occurred at this
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Fig. 8. (a) Fig. 7i which contains the non-migrated interface segments; (b) HRTEM image of the non-migrated interface segment region circled by red in (a); (c) magnification of the region marked by black rectangle shown in Fig. 7e; (d) HRTEM image of the migrated interface segment region circled by red in (c). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
interface. Fig. 9 shows the STEM-EDS results of the specimen derived at position 5. It is found that no large oxide particle is visible on the Cu/Cu-10Zn interface and all the grain boundaries initially on the Cu/Cu-10Zn interface have migrated at this position. Fig. 9e shows the magnification of the rectangle drawn in Fig. 9a. It is noted that a small oxide particle with a diameter of approximately 60 nm is sporadically visible in the grain interior. 3.2. Discussion The results obtained from Fig. 3 revealed that the initial Cu/Cu10Zn interface is completely unbonded and lots of oxides are densely distributed at the interface at position 1, i.e., the initial plastic deformation stage of the interface materials around the tool during the FSW. It is not surprising that the oxide regions in Fig. 3 show a higher O contrast in the O mappings and the comparatively lower Cu contrast in the Cu mappings because even the Cu-10Zn region has a high Cu content of ~90 at%. However, in the Zn mappings, the oxide regions show a much lower Zn contrast compared to the Cu-10Zn region. These results indicate that the oxides initially distributed on the Cu/Cu-10Zn interface at position 1 are mainly Cu oxides. However, when the interface evolved to position 2, i.e., an early plastic deformation stage of the interface materials, the small oxide particles distributed along the interface are seen to show the highest Zn contrast in the Zn mappings. All the oxide particles observed on the interface at the positions following position 2 are found to show similar characteristics. In Fig. 7i at position 4, an oxide particle detected on the interface marked by the yellow dotted circle is identified to be composed of ~29.1 at% O, ~42.8 at% Cu, and ~28.1 at% Zn based on the STEM-EDS spot analysis. Therefore, it is concluded that the oxide particles observed at
the interface positions after position 2 are likely a mixture of Cu oxides and Zn oxides. It is well known that the Cu oxides can react with Zn to form Zn oxides and Cu at high temperature because of the different reactivities of Cu and Zn, accompanied by a release of heat. Thus, it is inferred that during the interface material flow deformation around the tool in the FSW, the Cu oxides, which initially formed on the interface, reacted with Zn at an early plastic deformation stage, then a mixed oxide consisting of Cu oxide and Zn oxide finally formed on the interface with a reaction heat release. These findings might provide a new insight into taking advantage of the chemical reaction associated with an endothermic or exothermic behavior that occurred on the joining interface to possibly control the microstructure of the joining interface. The occurrence of the strain-induced grain boundary migration, which was pinned by the oxide particles, has been frequently identified on the Cu/Cu-10Zn joining interface during the FSW. The driving force for the boundary migration is regarded as the stored energy difference correlated to the dislocation density difference between the opposite sides of the boundary. Fig. 5c shows a schematic model for the strain-induced grain boundary migration [29]. Grain A and grain B have different dislocation densities, rA and rB , respectively. The DE, which is the stored energy difference between the opposite sides of the grain boundary, can thus be calculated according to the following equation: DE ¼ 12 GðrB rA Þb2 , where G is the shear modulus of the material, and b is the burgers vector of a dislocation. The bulging boundary shows a spherical cap shape with a radius of R and has a specific grain boundary surface energy of d. The condition for the grain boundary migration to occur is the favorable energy balance between the decrease in the total stored energy due to dislocation elimination and rearrangement caused by the passing boundary and the increase in total grain boundary surface energy due to the newly bulged boundary. This occurrence
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Fig. 9. STEM-EDS results of the foil specimen derived at position 5: (a) STEM bright field image of the specimen; (bed) corresponding Cu, O, Zn elemental mappings, respectively; (e) magnification of the region marked by the rectangle shown in (a); (feh) corresponding Cu, O, Zn elemental mappings of this region, respectively.
condition is thus given by the following formula: DE > 2Ld, where L is half of the distance between the intersections of the bulging boundary and the parent boundary. In the present study, since the strain-induced grain boundary migration is generally pinned by the distributed oxide particles, L can be regarded as half of the distance between the neighboring oxide particles. Therefore, it is concluded that the occurrence of the strain-induced grain boundary migration on the Cu/Cu-10Zn joining interface in the present study can be influenced by the following three factors, the stored energy difference between the opposite sides of the grain boundary, the specific grain boundary surface energy, and the interspacing between the oxide particles distributed on the interface. This view is well supported by the EBSD results shown in Fig. 10 which exhibits the Cu/Cu-10Zn joining interface microstructure derived from the location near position 3. The joining interface can be identified according to the distribution of the oxide particles indicated by the red arrows in the IQ map (Fig. 10b) and marked as the red dotted line in the IPF map (Fig. 10a). The interface regions indicated by the rectangles A, B, and C have different morphologies and distribution of the oxide particles thereby showing different interface behaviors. In region A, since almost no oxide particle is visible on the interface, the grain boundary migration occurs at this location; while for region B, no grain boundary migration is identified because lots of relatively large oxide particles are densely distributed on the interface leading to a small interspacing between the oxide particles. It is worth noting that even though the oxide particles are observed to be dispersed on the interface in region C, since the size and number density of the oxide particles are relatively small, the grain boundary migration can still occur at this position. In addition, it has been noted that the void is visible neighboring oxide particle 1, whereas no void is detectable beside oxide particle 2 according to Fig. 4eeh. Furthermore, it can also be seen that the strain-induced grain boundary migration occurring between these two oxide particles has overcome the pinning of oxide particle 2 but
still been pinned by oxide particle 1. These results suggest that the voids could form on the joining interface neighboring the oxide particles during the FSW because the existence of the oxide particles impedes the complete and ideal adhesion of the interface. The strain-induced grain boundary migration can promote the atomic diffusion and dislocation movement, rearrangement and annihilation in the nearby of the boundary, which will help to extinguish the retained voids adjacent to the oxide particles. Therefore, when the strain-induced grain boundary migration overcomes the pinning effect and passes through the entire oxide particles, the voids retained in the surroundings of the oxide particles are likely to be extinguished completely. It is also worth noting that the strain-induced grain boundary migration observed in the present study occurs from the Cu side towards the Cu-10Zn side. The reason for this phenomenon can be demonstrated as follows. Cu-10Zn is known to have a lower stacking fault energy of ~35 mJ/m2 compared to that of pure Cu (~78 mJ/m2) because of the alloying of Zn element [30]. Therefore, the full dislocations tend to dissociate into the partial dislocations and the dislocation cross-slip and dislocation climb become more difficult on the Cu-10Zn side, which makes the dislocation recovery associated with the dislocation mobility, rearrangement and annihilation more difficult thereby promoting the dislocation accumulation in the Cu-10Zn side compared with the Cu side. Therefore, it is easy for the Cu-10Zn side to exhibit a higher dislocation density compared to the Cu side, which causes the strain-induced grain boundary migration to easily occur from the Cu side towards the Cu-10Zn side due to the favorable stored energy balance. In the present study, the stacking fault energy difference between two work-pieces plays an important role on producing the stored energy difference between the opposite sides of the joining interface. However, it should be noted that even under the lack of stacking fault energy difference, the stored energy difference will also generate because the different grain orientations and different stress conditions between the opposite sides of the joining
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Fig. 10. EBSD results derived from the location near position 3.
interface can provide different dislocation generation rates and accumulation rates. In addition, it is seen that the apparent diffusion of Zn element across the Cu/Cu-10Zn interface can be identified in both the locations indicated by the red arrows in Fig. 7e and i based on the EDS line analysis of Fig. 7ee1 and i-1, and the location in Fig. 7e shows a relatively longer Zn diffusion distance of ~0.4 mm compared to that of the location in Fig. 7i (~0.25 mm). This difference in the Zn diffusion distance is considered to be associated with the fresh surface and atomic-level interface formation caused by the intense shear and compression as well as the thermal cycles and dwells at a sufficient temperature at these two locations during the FSW. In addition, it has been mentioned that the dislocations were strongly pinned and accumulated to form a highly-strained deformation zone at the interface in the location of Fig. 7e that facilitated the formation and growth of the dislocation-free nucleus inside it, while the interface in the location of Fig. 7i has not yet migrated. Both the formation of the highly-strained zone consisting of abundant defects and the formation and growth of the new grain accompanied with dislocation rearrangement and annihilation can promote the diffusion of the Zn element across the interface in the location of Fig. 7e, which can be one of the possible explanations for the relatively longer Zn diffusion distance in the location of Fig. 7e. Based on these results, the interface joining mechanism during the FSW can be concluded as illustrated in Fig. 11. For the initial joining interface of the work-pieces, it is completely unbonded and lots of large oxides are distributed along the initial joining interface as shown in Fig. 11a. The large oxides are then fragmented into small particles, and the oxide layers formed on the interface are ruptured due to the compressive force perpendicular to the interface and the shear force and interface material deformation parallel to the interface, which produces fresh surfaces on both Cu side and Cu-10Zn side as illustrated in Fig. 11b. Then an atomic-level joining is achieved between the fresh surfaces of both Cu and Cu-10Zn sides under the compressive force, while leaving the small oxide particles still dispersed on the interface with the voids formed around the oxide particles. As demonstrated in the magnified schematic inset in Fig. 11c, the dislocation slip and/or atomic diffusion occurring around the void assists to extinguish the void.
However, it is believed that a large deformation and/or high temperature is required to provide extensive dislocation gliding and/or abundant atomic diffusion around the oxide particles in order to completely extinguish the retained voids neighboring the oxide particles. For this interface condition, the stress concentration easily occurs at the locations of the oxide particles which may initiate the micro-cracks. The micro-cracks then easily propagate along the interface and connect to each other, then finally leading to a fracture along the interface. In fact, a strain-induced grain boundary migration will occur on the interface, which is driven by the stored energy difference associated with the dislocation density difference between the opposite sides of the grain boundary. This grain boundary migration will then promote the dislocation movement and atomic diffusion near the migrating boundary thus assisting to extinguish the voids formed around the oxide particles and also shift the locations of the oxide particles from the grain boundaries to grain interiors, as illustrated in Fig. 11d and e. Under further deformation, the oxide particles on the interface will be further fragmented into fine particles and further dispersed; finally, a sound interface joining without oxide particles and large voids remaining on the interface can be obtained during the FSW as demonstrated in Fig. 11f. 4. Summary In this study, a butt FSW process was performed using Cu and Cu-10Zn plates. Microstructural evolution along the Cu/Cu-10Zn interface in the fabricated FSW joint was systematically investigated to clarify the interface joining mechanism during the FSW. The obtained findings are shown as follows: (1) During the FSW, the large oxides initially existing on the work-pieces joining interface are fragmented into small oxide particles and the oxide layer is ruptured under the compressive force perpendicular to the interface and the shear force and interface material deformation parallel to the interface, which produces fresh surfaces on the interface. (2) The newly formed fresh surfaces of the work-pieces are then joined on an atomic-level under the compressive force, while
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Fig. 11. Schematic illustration of interface joining mechanism during FSW.
the small oxide particles are still dispersed on the interface with the voids formed around the particles. (3) A strain-induced grain boundary migration driven by the stored energy difference between the opposite sides of the boundary will actually occur at the interface, which facilitates extinguishing the voids formed around the oxide particles and shifts the locations of the oxide particles from the interface consisting of grain boundaries to the grain interiors. (4) The oxide particles are further fragmented and dispersed, and the strain-induced grain boundary migration is further developed under deformation, hence a sound interface joining without oxide particles and voids remaining on the interface is finally obtained during the FSW. Acknowledgements The authors wish to acknowledge the financial support by the New Energy and Industrial Technology Development Organization (NEDO) under the “Innovation Structural Materials Project (Future Pioneering Projects)”, JSPS KAKENHI Grant Numbers JP15H04133 and JP18K14027, and a ISIJ Research Promotion Grant. References [1] W.M. Thomas, E.D. Nicholas, J.C. Needham, M.G. Murch, P. Templesmith, C.J. Dawes, G.B. Patent Application No. 9125978.8 (1991). [2] R.S. Mishra, Z.Y. Ma, Friction stir welding and processing, Mater. Sci. Eng. R 50 (2005) 1e78. [3] C.G. Rhodes, M.W. Mahoney, W.H. Bingel, M. Calabrese, Fine-grain evolution in friction-stir processed 7050 aluminum, Scripta Mater. 48 (2003) 1451e1455. [4] K. Kitamura, H. Fujii, Y. Iwata, Y.S. Sun, Y. Morisada, Flexible control of the microstructure and mechanical properties of friction stir welded Ti-6Al-4V joints, Mater. Des. 46 (2013) 348e354. [5] S. Mironov, K. Inagaki, Y.S. Sato, H. Kokawa, Effect of welding temperature on
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