Insights into microstructure evolution and interface joining mechanism of ultrasonic spot welded Cu/Cu joint

To understand the underlying microstructure evolution and joining mechanism of ultrasonic spot welded Cu/Cu joint, the microstructure and crystallography analysis were investigated. The grain size, grain boundary, misorientation angle, Kernel Average Misorientation, Schmid factor and crystal orientation along the weld thickness were heterogeneous and unsymmetrical. The changes of corresponding microstructural and crystallographic features were more obvious for the top sheet comparing with the bottom sheet. The {111} <110> texture, nano-sized equiaxed induced by discontinuous dynamic recrystallization and refined elongated grains because of continuous dynamic recrystallization were formed at the welding interface. The nucleated fine grains between the elongated grains attributing to continuous dynamic recrystallization were beneficial to the grain boundary migration and consequent the achievement of the interfacial metallurgical bonding.


Introduction
Ultrasonic spot welding (USW) can produce coalescence through the simultaneous combined working of a high frequency tangential reciprocating mechanical vibration and applicable clamping pressure [1].In general, the formation process of a high quality joint includes three stages: slip stage, slip-stick transition stage and stick stage [2].Therefore, the relative motions at the faying interfaces of sonotrode tip/top sheet, top sheet/bottom sheet, and bottom sheet/anvil occur in the different welding phase, resulting in the heat generation induced by the friction, plastic deformation and the dissipation of ultrasonic energy.In addition, the materials at the welding interface bear a much high-strain rate (about 10 3 s −1 ) in less than one second [3].But the joining mechanism at the welding interface is not clear.
Recently, some researchers have studied the joining mechanism between two Cu sheets.It is reported that the plastic deformation first starts at the area beneath the peak of the sonotrode tip, and then spreads to the area beneath the valley of the sonotrode tip [4].The bond density and interlocking at the welding interface are enhanced with the increase of the welding time.Due to the limitations of the optical microscope, the accurate grain size, especially at the welding interface, and bonding formation mechanism are unclear.To resolve this issue, abnormal dislocation substructures and crystal orientation of the bonding interface for the Cu single crystals joint were studied in the reference [1].But the welding energy is rather low (only 200 J), thus the microstructure evolution result may be not in accordance with that under the actual working conditions.Therefore, large welding energy is desired.As the welding energy ranging from 400 to 2400 J increases, the welding interface with wavy pattern characteristic gradually takes place, and the low-angle boundary (LAGB) number first increases and then decreases [5].However, the quantified analysis of the crystal orientation and grain boundary character are not presented.Although the formation mechanism of the metallurgical bonding of Cu terminals to Cu cables were studied based on the crystal orientation and grain boundary character [6], the resolution ratio of electron backscatter diffraction (EBSD) is low, the grain size, grain boundary and its crystal orientation at the region that undergone severe plastic deformation are not achieved.Thus the formation mechanism at the corresponding area should be further clarified.In addition, the authors still didn't clearly illustrate the detailed dynamic recrystallization mechanism at the interface region, and how the boundary migration is done.
In our previous work [7], the interfacial bonding mechanism of Cu/Cu joint was investigated employing the transmission electron microscopy (TEM).The microstructure at the welding interface is inhomogeneous because of the complex welding process.Meanwhile, the area of TEM observation is relatively small, thus it does not prove a statistical significance.Hence, there were inevitably some limitations in the results.In comparison with the TEM, EBSD has merits of larger observation region and the collection of more statistically significant data on the grain boundary character, grain size, crystal orientation, etc.Thus systematical investigation on the joining formation mechanism at the Cu/Cu weld interface is desired based on the highresolution EBSD detection results.Moreover, almost all the previous works on USW is focused on the typical welding interface.It is envisioned that, because of the inherent nature of USW [2,3], a wide range of microstructure distribution along the thickness of weld zone is expected.Note that the research in this field has not been carried out yet.Therefore, in order to significantly advance our comprehending of USW, the current study focuses on understanding the evolution of the microstructure along the weld thickness and shedding light on the fundamental mechanism governing interface joining of USWed Cu/Cu joint.

Materials and experiments
The materials utilized in this study are commercial T2 Cu sheets, and the dimensions of the specimens are 0.5 and 1.0 mm thickness × 100 mm length × 25 mm width, respectively.Before welding, the surfaces of the raw Cu sheets were treated using 800 # sand paper, and then cleaned by the alcohol.Three-dimensional morphology and surface roughness of the treated Cu sheets, measured by confocal microscopy (Leica DCM8), are shown in Figure 1(a, b).There are numerous micro protrusions on the base metal surface.The surface roughness (Ra) is about 0.56 μm.A BWX-D2042 ultrasonic welder was employed to weld the Cu sheets.The 0.5 mm thick Cu sheet as the top specimen was presented directly beneath the sonotrode tip, and the 1.0 mm Cu sheet was considered as bottom sheet.Schematic illustration of the USW process is shown in Figure 2(a).The ultrasonic vibration direction is perpendicular to rolling direction (RD), and parallel to the transverse direction (TD).The clamping force is exerted on the top Cu sheet, and its direction is parallel to the normal direction (ND).The sonotrode tip size is 7 × 8.8 mm 2 .The detailed knurl patterns and corresponding dimensions on the surfaces of the sonotrode tip and anvil can be found in the reference [7].The using welding process parameters are 1800 N clamping force, 20 kHz frequency, 36 μm vibration amplitude and 0.8 s welding time.The selection of the parameters mentioned above depends on our previous research result [7].Under this condition, Cu/Cu joint with the optimal mechanical property was obtained.Figure 2(b) exhibits the typical optical microscope Cu/Cu joint cross-section.
From the previous study result [7], the interface bonding in the peak region is important to the joint quality.Thus the microstructure evolution and interface joining mechanism in the peak region is the focus of this study.
The cross-section of the Cu/Cu joint was observed by optical microscope (Olympus GX53).The microstructure characterization was performed by scanning electron microscopy (SEM, FEI Verios 460) possessing the EBSD system.The grain boundarry with misorientation angle in the range of 2°to 15°were considered as LAGB and those with more than 15°were defined as high-angle boundary (HAGB).The boundary with 3 feature was defined as twinning boundary (TB).The data were analyzed using OIM software.

Microstructure characterization along the Cu/Cu weld thickness
Figure 3 shows the IQ (image quality) + grain boundary images, IPF (inverse pole figure) image, KAM (Kernel Average Misorientation) image and SF (Schmid factor) image of the Cu/Cu joint cross-section.From Figure 3(a), the bonding interface is demonstrated by the black arrows.The black region at the welding interface is not resolved, attributing to the existence of the unbonded gaps.In addition, there are obvious differences in the distributions of the grain boundary and size, crystal orientation, KAM and SF from the top specimen to the bottom specimen (Figure 3, Figures S1-S6).It is demonstrated that the materials along the weld thickness may undergo the different fields of the temperature, stress, stress strain and ultrasonic vibration.In order to rationalize more clearly how the microstructure evolves, the data are divided into five regions, as exhibited in Figure 3(b).
The average sizes of grain in the regions I, II, III, IV and V are 7.18, 5.77, 4.05, 14.65 and 22.55 μm, respectively (Figure 4(a)).Comparing with the top sheet (0.5 mm), more grains with smaller size are existed in the regions I, II and III, while grains with larger size decrease significantly; and the average sizes of grains in the regions I, II and III are reduced by 25.4%, 40.0% and 57.9%, respectively.Comparing with the bottom sheet (1.0 mm), more grains with smaller size are present in the region III, and the average size of grain is reduced by 46.9%.However, more grains with larger size of grain are observed in the regions IV and V, and the average size of grain is increased by 92.0% and 195.5% respectively.It is demonstrated that the grain in the regions I, II and III are refined, but the grain growth occurs in the regions IV and V.
The average misorientation angle in the regions I, II, III, IV and V are 35.0°,30.9°, 33.2°, 45.5°and 46.1°, respectively (Figure 4(b)).In comparison with  the corresponding original Cu sheet, numerous grain boundaries possessing the misorientation angles lower than 55°are observed, while the grain boundaries having the misorientation angles higher than 55°reduce obviously in all the five regions.The corresponding detailed fractions of the LAGB, HAGB and TB are shown in Figure 4(c).In comparison with the top sheet, the LAGB fractions in the regions I, II and III are increased by 88.51%, 149.32% and 62.16%, respectively; the HAGB fractions are decreased by 15.38%, 25.94% and 10.80%, respectively; and the fractions of the TB are reduced by 53.21%, 67.11% and 79.14%, respectively.In contrast to the bottom sheet, the LAGB fractions in the regions III, IV and V are enhanced by 126.42%, −15.09% and −29.25%, respectively; the fractions of the HAGB are decreased by -14.99%, 1.79% and 3.14%, respectively; and the TB fractions are reduced by 78.69%, 41.80% and 21.58%, respectively.This trend can prove that the degree of plastic deformation along the weld thickness is not uniform.
The average KAM in the regions I, II, III, IV and V are 1.02°, 1.36°, 0.64°, 0.37°and 0.23°, respectively (Figure 4(d)).Comparing with the top sheet, higher KAM in the regions I, II and III are present, and the average KAM are enhanced by 85.5%, 147.3% and 16.4%, respectively.In contrast to the bottom sheet, higher KAM in the regions III are found, and the average KAM is increased 48.8%; but lower KAM in the regions IV and V are detected, and the average KAM are reduced by 14.0% and 46.5%, respectively.
The average SF in the regions I, II, III, IV and V are 0.445, 0.454, 0.481, 0.425 and 0.432, respectively (Figure 4(e)).Comparing with the top sheet, higher SF in the regions I, II and III are observed, and the average SF are increased by 0.5%, 2.5% and 8.6%, respectively.In contrast to the top sheet, higher SF in the regions III and V are existed, and the average SF are increased by 12.4% and 0.9%, respectively; but lower SF in the region IV are detected, and the average SF in the region IV is reduced by 0.7%.
During USW, the sonotrode tip directly acts on the top sheet (Figure 2), thus more friction heat and ultrasonic softening affect are generated on the top sheet [8].Moreover, the clamping force directly acts on the top sheet, leading to the higher compressive stress in comparison with the bottom sheet.Therefore, the plastic deformation on the top sheet is much severer than that on the bottom sheet, which in turn promotes the formation of the plastic deformation heat.During the severe plastic deformation, dislocation slip is the main deformation mechanism.Amounts of dislocations are produced in the grain.Dislocation walls are formed through the slip accumulation and entanglement of dislocations.As the degree of plastic deformation increases, the dislocation density increases, and the dislocation walls change to the LAGB and further evolve into the HAGB [9].Then the coarse grain is divided into the small size grains.As a result, the grain refinement, the increasing LAGB fraction and the decreasing HAGB fraction are present in Figures 3(a)  and 4(a, c).The high density of the LAGB is an evidence of recovery of dislocations into sub-grain boundaries.TB can resist the dislocation slip transmission, but the resistance degree depends on the dislocation type [10].Due to the severe plastic deformation occurrence, more dislocations will be incorporated into more TBs, leading to the structural changes of the corresponding TB.Consequently, the TB fraction decreases ((Figures 3(a) and 4(b, c)).Interestingly, the relative movement with the high frequency and large amplitude also take place at the welding interface, and the shear ultrasonic vibration and the dynamic shear stress are formed, promoting the further plastic deformation generation.The plastic deformation at the welding interface continues to increase, accelerating to the further grain refinement.As a consequence, the grain size from the region I to III decreases (Figures 3(a) and 4(a)).The maximum shear stress is at the sliding surface for the both Cu sheets.The largest degree of plastic deformation must have occurred at the welding interface (Figures 3(a, d) and 4(e)).Due to the damping of the shear stress and ultrasonic softening effect, a gradient plastic deformation occurs (Figure 3).In theory, the LAGB fraction and the dislocation density should be the largest at the welding interface (Region III), but it is indeed not the case.This is because of the formation of the dynamic recrystallization, and its formation mechanism will be systematically elaborated in the next section.Thus the LAGB fraction and the average KAM angle in the region II are larger than that in the regions I and III (Figures 3(a, c) and 4(b-d)).
Because of the synergistic action of the lower friction and plastic deformation heat, small compressive stress and weak ultrasonic softening effect on the bottom to the incomplete dynamic recrystallization.The average sizes of the equiaxed grain in the regions IV and V are 92.0%and 195.5% larger than that in the bottom sheet, respectively (Figures 3(a) and 4(a)).It is demonstrated that the grains are noticeably grown up.The reasons are exhibited as below.During USW, the heat is mainly transmitted from the top sheet to the bottom Cu sheet [11].Meanwhile, the friction energy also is generated at the bottom sheet/anvil faying interface due to the relative movement at the corresponding interface.It is inferred that the absorbed energy in the region IV is more than that in the region V.The more energy the materials absorb, the higher temperature the region is.The driving force of the grain growth is to reduce the total interfacial energy.In fact, as the grain grows, the grain boundaries always moves towards the curvature center and are gradually straightened.So the small grains disappear.Higher temperature can accelerate the grain growth and dislocation annihilation.Thus the region V has the larger size of grain, the higher average misorientation angle and HAGB fraction, and the lower LAGB fraction and KAM in comparison with that in the region IV and the 1.0 mm Cu sheet.
Figure 5 exhibits IPF images of base metals and Cu/Cu joint.In the case of the IPF images obtained from specimens along TD, the main textures of < 111 > / < 101 > and < 111 > / < 001 > are presented in the top and bottom sheets, respectively.In comparison with the top sheet, the < 111 > direction cannot be observed, the < 101 > texture weakens and the < 212 > texture occurs in the region I.For the regions II and III, no < 212 > texture exists, but the aggregation of orientation near the < 101 > direction can be observed, and the corresponding intensity increases.It is concluded that the < 111 > direction of the top sheet first translates into the < 212 > direction in the region I, then the < 101 > direction.Its intensity gradually increases from the region I to III.In contrast to the bottom sheet, the < 111 > direction weakens step by step, while the density of the < 001 > direction is enhanced in the regions IV and V.
In the case of the IPF images achieved from samples along ND, the dominating textures of < 101 > and < 101 > / < 112 > are presented in the top and bottom sheets, respectively.In contrast to the top sheet, the < 101 > direction weakens, and the < 213 > direction is observed in the region I.The < 101 > direction almost disappears and the < 112 > direction is formed in the region II.For the region III, the dominating direction is < 111 > .It is demonstrated that the < 101 > direction first changes to the < 213 > direction, then the < 112 > direction and last the < 111 > direction.Comparing with the bottom sheet, the < 101 > / < 112 > directions weaken, and the < 001 > direction is detected in the region IV.The density of the < 001 > direction is improved, but the density of the < 101 > / < 112 > directions consistently decrease in the region V.
In the case of the IPF images observed from RD, the major directions of < 001 > / < 111 > and < 114 > / < 101 > are presented in the top and bottom sheets, respectively.In contrast to the top sheet, the < 111 > direction weakens step and step, and the < 001 > direction intensity increases in the regions I, IIand III.Comparing with the bottom sheet, the weak < 114 > direction almost disappears and the < 101 > direction density decreases continuously, but the density of the < 001 > direction is gradually enhanced in the regions IV and V.In short, for the IPF distributions of RD direction, all the other directions weaken and eventually change to the < 001 > direction.Overall, the rotation behaviors of the grains are heterogeneous along the weld thickness.
The dislocation activities are critical to the texture transition, because texture formation is always dislocation-based [12].The correlation between texture transition and dislocation glide is build by detecting the SF of all the grains.When the resolved shear stress arrives to the critical value, the dislocation slip will occur in the slip system possessing the feature of the higher SF.Generally, {111} < 110 > slip system is critical to the plastic deformation in pure copper.For the regions I, II and III, in comparison with the top sheet, the average SF are increased 0.5%, 2.5% and 8.6%, respectively (Figures 3(d) and 4(e)).The damping of the shear stress and ultrasonic softening effect occurs from the region III to I. It is demonstrated that larger shear stress and ultrasonic vibration can accelerate the lattice rotation towards the easier slip condition, leading to the drop of the flow stress [13,14].During plastic deformation, the grain orientation will change from the minimum angle to the maximum angle, when the grain is rotated to an appropriate direction where the resolved shear stress on this activated slip system can reach the maximum value, the other slip systems will be triggered [15].Thus, seen from TD, the orientation directions ( < 111 > , < 001 > ) with lower SF in the top and bottom sheets gradually turns into the < 101 > direction (Figures 3(b) and 5) with higher SF (Figures 3(d) and 4(e)) in the region III.Obtained from ND, the < 101 > direction in the both sheets gradually changes to the < 111 > direction in the region III (Figure 5, Figure S5(b)).This indicates that {111} plane is parallel to the RD-TD plane.Consequently, the {111} < 110 > texture is formed at the welding interface (region III).In turn, it is beneficial to plastic deformation.The variation of the texture seen from RD (Figure 5, Figure S5(a)) is induced by the lattice rotation along the TD and ND directions.For the regions IV and V, it is thought that the surface energy minimization could lead to a texture formation when the grains grow.The grains having the < 111 > direction in pure Cu possess the lower surface energy [16].But the < 111 > orientation density noticeably decreases, while the < 001 > orientation intensity increases in comparison with the bottom sheet (Figure 5).It is in accordance with the previous result [12].

Interface joining mechanism
From Figure 6(a), the welding interface is demonstrated using the black arrows.It can demonstrate that the both base metals are successfully welded.A thin region that  and then the elongated cells are break down and replaced by the small elongated grains (Figure 6(a,  b)), with the decreasing spacing and the relatively low KAM (Figure 6(c)).The reasons can be summarized as below.During USW, the maximum shear stress is at the faying surface, and decreases towards to the bulk material.It is rational that at the center of the welding zone, dislocation movement must have been much more pronounced, leading to the increased shear degree.This can be demonstrated by the obvious orientation (Figure 6(b)) and SF (Figure 6(d)) variations at the welding interface in comparison to the base metal.In theory, the shear degree along the welding interface towards to top and bottom sheets should be symmetrical in that sense.But that's not what the experiment results show.In comparison with the bottom sheet, more friction and plastic deformation heat and ultrasonic softening effects are acted on the top sheet [8], leading to the obvious decrease of the yield strength.Consequently, under the same conditions, e.g. the same shear stress and ultrasonic vibration, Cu at the top sheet side is more deformed than that at the bottom sheet side at the welding interface.Thus the welding zone is a thin area that extends different distances into the both sheets from the welding interface.Meanwhile, the existence of LAGBs and new HAGBs indicate that the deformation mechanism is the slip of dislocations.The gradual and continuous transition from the inner HAGBs to the outer LAGBs demonstrates the formation of the continuous dynamic recrystallization at the welding zone and the dynamic recovery at the periphery of the welding zone.For the welding zone, the newly formed small grains, < 101 > direction and the high value of the HAGB fraction and SF demonstrate the generation of the large lattice rotations during continuous dynamic recrystallization, due to the sliding friction and frictional heat.
The nucleation of dynamic recrystallization preferentially occurs at the sites of high dislocation density.The nucleated fine grains between the elongated grains are observed at the bulging grain boundary, as shown in the regions E, F and G from Figure 6(b).Under the combined action of the shear stress and the elevated temperature, the displacement of grain boundaries parallel to the shear direction will take place.It will promote the lattice rotation of the bulging part under the reciprocating shear plastic deformation [17].The more reciprocating cycles can accelerate the formation of the larger misorientation and consequent the separation of a new grain from the parent grain, e.g. the regions E, F and G in Figure 6(b).The grains indicated by the black arrows mark in the region F from Figure 6(b) is such a dynamic recrystallization grain just formed by the lattice rotation.This grain formation process is also in accordance with the feature of the continuous dynamic recrystallization.It is indicated that continuous dynamic recrystallization is the main mechanism for the grain occurrence between the elongated grains.This is beneficial to the grain boundary migration and consequent the achievement of interfacial metallurgical bonding.As the strain rate increases, the preference of dynamic recrystallization nucleation at the grain boundary gradually disappears [17].The trend of the grain size in the regions from C to F is increasing.It is indicated that the strain rate decreases from the weld interface to the base metal.
Note that in some regions, e.g. the locations A and B in Figure 6(a), nano-sized equiaxed grains with random orientation (Figure 6  stage, the amplitude of the relative motion between the both sheets is large [1].The micro asperities (Figure 1) at the original Cu sheets' surfaces first touch each other, thus the intense plastic deformation with high strain rate ( ∼ 10 3 s −1 )[3] is generated, resulting in the dislocation formation.High density dislocations can promote the generation of the large nucleation rate and the small recrystallized grain size.When the thermal activation energy induced by the increased welding interface temperature due to the friction and plastic deformation heat [8] as well as the stored energy because of the dislocation multiplication exceed the threshold, the discontinuous dynamic recrystallization occurs.As a result, in some regions, e.g. the locations A and B in Figure 6(a), the random orientation, high HAGB fraction and low KAM are present in the nano-sized equiaxed grains.The suppression of the nano-sized equiaxed grain growth is attributed to the low welding temperature and its rapid cooling speed [7].Interestingly, SF in one grain indicated by the black cycle at the weld zone along the bottom sheet side is very small (Figure 6(d)).It is illustrated that slight plastic deformation takes place at the welding interface along the bottom Cu sheet side.

Lap shear strength
Figure 7 shows the typical lap shear stress-strain curves of the Cu/Cu joints and the failure modes (see insert).The lap shear strength is 37.8 MPa.It is indicated that the bonded areas at the welding interface can deliver the heavy load during the loading process.It could be ascribed to the synergistic effects of the interface metallurgical bonding and the refined grains.

Conclusions
(1) The grain boundary, grain size, KAM, misorientation angle, SF and crystal orientation along the weld thickness are heterogeneously unsymmetrically distributed.The changes of corresponding microstructural and crystallographic features are more obvious for the top sheet in comparison with the bottom sheet.Slight plastic deformation takes place at the welding interface along the bottom Cu sheet side.(2) Due to the severe plastic deformation, the {111} < 110 > texture, the nano-sized equiaxed induced by the discontinuous dynamic recrystallization, the refined elongated grains because of the continuous dynamic recrystallization, the decreased fraction of TB and the increased SF are formed at the welding interface.Because of the formation of the dynamic recrystallization, the high HAGB fraction, low KAM, small grani size are generated at the welding interface.(3) The nucleated fine grains between the elongated grains due to continuous dynamic recrystallization can promote the grain boundary migration, leading to the formation of the interfacial metallurgical bonding.The strain rate decreases from the weld interface to the base metal.

Figure 2 .
Figure 2. (a) Schematic illustration of USW, (b) the typical optical microscope of Cu/Cu joint cross-section.

Figure 3 .
Figure 3. (a) IQ + grain boundary images, (b) IPF image, (c) KAM image and (d) SF image of the cross section of Cu/Cu joint.

Figure 5 .
Figure 5. IPF images of base metals and Cu/Cu joint.

Figure 6 .
Figure 6.(a) IQ + grain boundary images, (b) IPF image, (c) KAM image and (d) SF image from the high-resolution EBSD of the welding interface.
(b)) are observed, meanwhile almost no LAGBs (Figure 6(c)) are observed.The reasons can be summarized below.At the initial USW

Figure 7 .
Figure 7.Typical lap shear stress-strain curves of the Cu/Cu joints and the failure modes (see insert).