2023 Volume 64 Issue 4 Pages 827-834
Cu/A5052 dissimilar metal laminates (DMLs) were fabricated by an accumulative roll bonding process. The Cu/A5052 DML consisted of hard and soft layers, in other words, high and low shear modulus layers, which is called a mille-feuille structure (MFS). The Cu/A5052 DML with the MFS showed kink deformation during monotonic compression. Kink formation was analyzed by in-situ optical microscopy and digital image correlation. The development behavior and morphology of the ortho-type kink depended on the annealing conditions prior to the compression tests. Based on the observations and measurements, criteria for kink formation in the DML with the MFS were proposed.
Fig. 4 Snapshot images, shear strain maps, and compressive strain maps of the ARB-processed Cu/A5052 composite during compression tests evaluated by DIC. The left column is for As-ARB, and the right column is for PA573. The nominal compressive strains of (a) and (g), (b) and (h), (c) and (i), (d) and (j), (e) and (k), and (f) and (l) are 0.04, 0.08, 0.12, 0.16, 0.20, and 0.24, respectively.
Kink deformation during compression tests has been observed in various materials, such as intermetallic compounds,1) polymer-matrix/carbon-fiber composites,2) abalone shells consisting of brittle and ductile layers,3) ice,4) and even strata, the scale of which is on the order of kilometers,5) since Orowan reported kink deformation in cadmium in 1942.6) Kink deformation has also been found in metallic alloys. In particular, some Mg alloys with a long-period stacking order (LPSO) microstructure7,8) are known to show clear kink deformation.9–14) One of the most notable mechanical properties of LPSO-Mg alloys is that the alloys are significantly strengthened by introducing deformed kinks,15–17) which was originally reported in a Mg–Zn–Y alloy by Kawamura et al.18) Therefore, kink introduction due to deformation in materials is highly expected to be a new strengthening method from an engineering point of view. Recent studies have revealed the three-dimensional morphology of kink structures both experimentally19) and theoretically.20)
The kink-associated deformation behavior in LPSO-Mg alloys is attributed to the limited number of slip systems in hexagonal close packed (HCP) crystals, i.e., the slip plane is limited to the basal plane at room temperature. However, previous studies have not sufficiently revealed the yielding phenomenon related to kink formation in LPSO-Mg alloys. There has also been no quantitative theoretical explanation of the yield stress, which is strongly affected by kink formation. Furthermore, the morphology of deformed kinks in LPSO-Mg alloys may be different from those reported in other metallic materials.
Apart from LPSO-Mg alloys, Nizolek et al. showed kink formation in Nb–Cu dissimilar metal laminates (DMLs) fabricated by an accumulative roll bonding (ARB) process.21) Hagihara et al. also recently reported kink formation in Al/Al2Cu eutectic alloys with an alternating layered microstructure.22) Such metallic materials consisting of elastically soft and hard layers showing kink formation have been called “mille-feuille structure” (MFS) materials.22) LPSO-Mg alloys can be said to be typical MFS materials.
To understand the kink deformation of metallic materials with the MFS, another multilayer composite different from the LPSO phase is required. DMLs are thought to be a model material for studying kink deformation. Recently, criteria for obtaining a satisfactory layered structure using the ARB process were proposed,23) and they are applicable to fabricating Cu/A5052 DMLs with the MFS. In this study, therefore, Cu/A5052 DMLs were fabricated by the ARB process as the model material with the MFS, and then, the kink formation was investigated through compression tests at room temperature. Furthermore, the effects of preannealing before the compression tests on the morphology of developed kinks were also examined in terms of favorable conditions to introduce deformed kinks.
Pure Cu sheets with a thickness of 1 mm and A5052 Al alloy sheets with a thickness of 0.5 mm were subjected to the ARB process for this study. The purity of the Cu sheets was 99.98 mass%, and the chemical composition of the A5052 sheets is shown in Table 1. The Cu and A5052 sheets for the ARB process were previously annealed at 873 K and 673 K, respectively, for 7.2 ks under ambient conditions using an electric furnace (KBF828N1, Koyo Thermo). Cu/A5052 DMLs were fabricated by adopting the ARB process, and a schematic diagram is shown in Fig. 1(a).23) The coordinate system of a sample is defined by the rolling direction (RD), transverse direction (TD), and normal direction (ND). Planes with normal direction parallel to the RD, TD, and ND are denoted as the RD, TD, and ND planes, respectively. Hereafter, a specimen for which the ARB process has been applied n times is denoted ARB nc. The as-annealed sheet without application of the ARB process is denoted ARB 0c.
Schematic illustration of the (a) ARB process and (b) compression tests recorded with a camera.
As the first ARB process, a Cu sheet was sandwiched between two A5052 sheets, after both surfaces of the Cu sheet and the facing sides of the A5052 sheets were cleaned by acetone and treated by wire brushing to remove the surface contamination and oxidized layers. Cold rolling (CR) with a rolling reduction of 50% was applied in the first ARB cycle, followed immediately by water quenching. After the first ARB cycle, a standard ARB process was conducted, i.e., the sheets were cut into two sheets in the RD, and surface treatments, stacking, and roll bonding were applied. The details of the standard ARB process can be found elsewhere.24) Practically, the maximum number of cycles is limited due to the appearance of cracks from the sides of the sheet during CR. In this study, the ARB process was performed up to 7c. Intermediate annealing at 573 K for 1.8 ks was subsequently applied to the sheet after the 3rd and 6th ARB processes to reduce the strength of each layer of the Cu/A5052 DMLs based on the criteria to maintain the parallel-multilayer structure without necking defects.23)
Rectangular shaped specimens were cut using an arc discharge wire cutting machine (HS-300, Brother) for the compression tests. The size of the specimens was approximately 2 mm (RD) × 1.1 mm (ND) × 2.5 mm (TD). Prior to compression tests for kink formation, a preannealing process at either 573 K or 673 K for 3.6 ks was applied for some of the ARB 7c sheets to reduce the strength of each layer. Hereafter, the as-ARB-processed sample is referred to as As-ARB. Additionally, the samples preannealed at 573 K and 673 K for 3.6 ks prior to compression tests are referred to as PA573 and PA673, respectively.
2.2 Microstructure observationsBackscattered electron (BSE) imaging and electron backscattering diffraction (EBSD) were applied on the TD plane of the specimens using a field emission type-scanning electron microscope (FE-SEM: JSM-9000F, JEOL). The acceleration voltage was 20 kV, and inverse pole figure (IPF) maps were constructed using analysis software (Aztec, Oxford Instruments). SEM specimens were cut by an electron discharge machine, and then, the surface of the TD plane was mechanically polished by emery paper up to #2000 to remove the damaged layer caused by the arc discharge. The surface of the TD plane was finally polished by an ion milling machine (IM4000PLUS, Hitachi High-Tech) in cross-sectional milling mode for 86.4 ks. The acceleration voltage and discharge voltage of the cross-sectional ion milling were set to 3 kV and 1.5 kV, respectively.
2.3 Chemical composition analysisEnergy dispersive X-ray spectroscopy (EDS) was performed on the TD plane of the specimens using an SEM (JSM-6390A, JEOL) with an energy dispersive X-ray spectrometer (JED-2300A, JEOL). The acceleration voltage was set to 20 kV. The TD plane of the specimens was mechanically polished by emery paper up to #4000 and finished by buffing with α-alumina with a diameter of 0.06 µm.
2.4 Compression testsCompression tests were conducted using a universal testing machine (AG-Xplus 10 kN, SHIMADZU) at room temperature (R.T.) in ambient conditions. The stress axis was parallel to the RD, and the initial strain rate ε0 was 10−3 s−1. A schematic illustration of the tests is shown in Fig. 1(b). In-situ optical microscopy was performed during compression tests on the TD plane using a high-speed camera (FASTCAM Mini WX100, Photron) with illumination by a light-emitting diode (LED). The TD plane was mechanically polished by emery paper up to #2000 before the tests. The optical images were collected with a time resolution of 1 s and analyzed by digital image correlation (DIC) using VIC-2D (Correlated Solutions). The image resolution of the camera was 2048 pixels × 2048 pixels, and the subset and step sizes were set to 11 pixels and 5 pixels, respectively.
Figure 2 shows nominal stress σn – nominal strain εn (s-s) curves from compression tests of As-ARB, PA573, and PA673. The black dotted lines in the s-s curves were drawn at strain intervals of 0.04, which show the timing of the images displayed in Fig. 4. For As-ARB and PA573, plastic deformation started at εn ≈ 0.06. The corresponding 0.2% proof stresses σ0.2 were 374 MPa and 307 MPa for As-ARB and PA573, respectively. The s-s curves of both specimens showed a plateau of approximately 420 MPa after yielding. Contrary to As-ARB and PA573, PA673 showed a much higher yield stress close to 800 MPa and fractured without any plastic deformation. PA673 showed brittle fracture at the point where σn dropped. The morphology of the brittle fracture will be discussed later.
Stress–strain curves from the compression tests of specimens As-ARB, PA573 and PA673 with an initial strain rate of 10−3 s−1.
Figure 3 shows BSE images of As-ARB on the TD plane under (a) As-ARB conditions and (b) after compression with εn over 0.3. Hereafter, x and y axis are defined being parallel to ND and RD, respectively. Figure 3(c) shows PA573 after compression with εn over 0.3. Figure 3(a) shows the laminated structure, in which the width of each layer was approximately 10 µm after ARB 7c, and the light and dark layers correspond to Cu and A5052, respectively. The Cu/A5052 DML was composed of alternating hard (Cu) and soft (A5052) layers. In other words, Young’s modulus of Cu is higher than that of A5052. Furthermore, the parallel layered structure was well maintained. As shown in Figs. 3(b) and (c), a kink (the region between red dashed lines) was clearly formed after the compression tests of As-ARB and PA573. More precise observations of kink formation during compression tests were conducted using the DIC technique as follows.
BSE images of the as-fabricated Cu/A5052 composite (As-ARB) (a) before and (b) after the compression test and of (c) specimen PA573 after compression. The red dashed lines represent the kink boundaries. It is noted that ND and RD are x and y axis, respectively.
Both As-ARB and PA573 showed kink deformation, and in-situ optical microscopy was applied during compression tests of both specimens to investigate the effects of preannealing on kink formation. Figure 4 shows the specimen images on the TD plane during compression tests, the recording timing of which is indicated by the dashed lines in Fig. 2. Figures 4(a)–(f) in the left column are for As-ARB, and Figs. 4(g)–(l) in the right column are for PA573. In each figure, the optical image captured by the camera is displayed on the leftmost side. The local shear strain εxy and the compressive strain εyy maps were constructed by applying DIC to the optical images, and both maps are displayed as color maps to the right of the original optical images. It is pointed out that subscript of x and y denote the x and y axis, respectively. Here, the range of local strain is between −0.2 and 0.2, as represented by the color bar, and the compression direction (CD) is parallel to the vertical direction in the images.
Snapshot images, shear strain maps, and compressive strain maps of the ARB-processed Cu/A5052 composite during compression tests evaluated by DIC. The left column is for As-ARB, and the right column is for PA573. The nominal compressive strains of (a) and (g), (b) and (h), (c) and (i), (d) and (j), (e) and (k), and (f) and (l) are 0.04, 0.08, 0.12, 0.16, 0.20, and 0.24, respectively.
First, the compression of As-ARB is considered using Figs. 4(a)–(f). In the initial stage of deformation (εn = 0.04 and 0.08), uniform deformation seemed to occur based on the optical images displayed in the leftmost column in Figs. 4(a) and (b). There was no localized strain, represented as no color gradation in either the εxy or εyy map. When εn = 0.12 (Fig. 4(c)), the optical image showed slight nonuniform deformation and barreling in the bottom half of the right side. Strain localization was not detected in the εxy map, whereas it was detected in the εyy map at the top-right corner. The localized strain appeared as a narrow band along the 45° direction with respect to the CD from the bottom-left corner. The narrow band seemed to stop within the specimen; in other words, it did not penetrate through the specimen. When εn = 0.16 (Fig. 4(d)), nonuniform deformation and barreling progressed simultaneously. Here, a clear dark band could be clearly seen along the 45° direction with respect to the CD from the bottom-left corner to approximately the halfway point of the right side. Such a band was also detected in both the εxy and εyy maps. The narrow band-like deformation remarkably progressed at εn = 0.20, as shown in Fig. 4(e). Furthermore, not only the bottom right of the specimen but also the left side of the specimen seemed to deform like a barrel. The band region became wider than that at εn = 0.16. The band region was also seen in both the εxy and εyy maps. As shown in Fig. 4(f), the band region propagated much wider and barreling progressed at εn = 0.24. Note that the top and bottom of the specimen were clearly horizontally misaligned at this stage. Both the εxy and εyy maps showed similar changes. Strictly speaking, the magnitude of εyy was always higher than that of εxy within the band. For instance, at εn = 0.24, the absolute value of εyy was higher than that of εxy. The region with higher εyy was larger than that with higher εxy. In other words, εyy was generated prior to εxy. Such localization of εyy could be a precursor phenomenon to trigger of εxy.
Next, the compression of PA573 is considered using Figs. 4(g)–(l). When εn reached 0.08 (Figs. 4(g) and (h)), the specimen seemed to show uniform deformation based on the optical images. Both the εxy and εyy maps also showed a uniform distribution. When εn reached 0.12 (Fig. 4(i)), the optical image and εxy map revealed that uniform deformation still occurred. However, an increase in εyy was detected at the bottom of the sample. Furthermore, the leftmost of the bottom region showed the highest absolute value of εyy. When εn reached 0.16 (Fig. 4(j)), a relatively clear dark line along the 45° direction with respect to the CD appeared from the lower left corner to approximately the halfway point of the right side of the specimen (see optical image). The εxy map also showed a narrow band-like region with a higher absolute value of εxy compared with the other regions. The εyy map showed a larger and unconstrained region with a higher absolute value of εyy compared with the other regions. Note that the highest absolute value of εxy was lower than that of εyy. We note that the band detected as a strain localized region in the εyy map was wider than that detected in both the optical image and εxy map. When εn reached 0.20 (Fig. 4(k)), the strain-localized region became wider, as certainly detected from the optical image and εxy map. Compared with As-ARB, i.e., Fig. 4(e), a larger and more localized displacement along the horizontal axis appeared at the end of the right-side band in Fig. 4(k) of PA573. The band-like region can also be recognized in the εxy map. The detected region with a high absolute value of εyy seemed slightly wider in the εyy map. At εn = 0.24 (Fig. 4(l)), the strain-localized band became wider, as shown in the optical micrograph and εxy map. In contrast, the band region was unclear in the εyy map, while the high strain region slightly expanded from the state in Fig. 4(k). Note that the size and morphology of the high-strain region detected in the εxy map did not completely coincide with those detected in the εyy map at any value of εn.
It is pointed out again that εn is 0.30 for Fig. 3, but, εn is up to 0.24 for Fig. 4. Thus, the strain localization area of As-ARB seems narrower than that of PA573 in Fig. 3. It is emphasized that Fig. 4 is displayed for discussing the initiation of the strain localization phenomenon during compression. Nevertheless, the formation of an ortho kink or a kink band was clearly observed by in-situ optical microscopy during compression tests and DIC analysis of the Cu/A5052 DML with the MFS in this study.
3.4 SEM/EDS and FE-SEM/EBSD measurementsBSE images and SEM/EDS line analyses of the ARB-processed Cu/A5052 composites are shown in Fig. 5. Figures 5(a) and (b) correspond to As-ARB, Figs. 5(c) and (d) to PA573, and Figs. 5(e) and (f) to PA673. The blue dots in Figs. 5(a), (c), and (d) represent the points for the SEM/EDS line analyses displayed in Figs. 5(b), (d) and (f), respectively. The BSE images clearly showed a distinct phase boundary in As-ARB (Fig. 5(b)), whereas an unclear narrow phase boundary appeared in PA573 (Fig. 5(c)). Furthermore, PA673 clearly showed some layers with the gradation of the contrast, as shown in Fig. 5(e). Such fine boundary interfaces without an intermetallic layer at the phase boundary observed in As-ARB were also reported when the ARB process was used to fabricate DMLs.23) In contrast, the unclear interface in PA573 and the gradation region in PA673 were considered to be intermetallic layers consisting of Al and Cu formed by diffusion during annealing.
BSE images and SEM/EDS line analyses of (a), (b) As-ARB, (c), (d) PA573, and (e), (f) PA673. The line analyses were performed at the points represented as dots in (a), (c) and (e).
The SEM/EDS line analyses also support the formation of a sharp interface without intermetallic layers in As-ARB (Fig. 5(b)), a thin intermetallic layer at the interphase interface in PA573 (Fig. 5(d)), and an alternating composition gradient at the interphase interface in PA673 (Fig. 5(f)). The SEM/EDS line analyses (Figs. 5(b) and (d)) indicated that the compositions of Cu within Al and Al within Cu adjacent to the interphase interfaces in As-ARB were lower than those in PA573. In other words, the width of each layer with 100% Cu and Al in As-ARB seemed wider than that in PA573. Figure 5(e) indicates that the composition of the gradation region cannot be considered to be similar to the composition of the original Cu and A5052 layers in PA673. The width of the 100% Cu layer in PA673 was much narrower than that in As-ARB and PA573. In the case of PA673, the width of the 100% Al layer was the narrowest among all the layers in all the specimens. As shown in Fig. 5(c)–(f), the thicknesses of the intermetallic layers of PA573 and PA673 were less than 1 µm and approximately 5 µm, respectively.
It can be said that the multiple types of intermetallic compounds are formed at the interphase-interface in PA673. The gradient in Fig. 5(f) also supports it. It is also pointed out that the special resolution of EDS measurement is limited since W-type SEM/EDS is used in this study. Therefore, the determination of types of intermetallic compounds was not performed.
Figure 6 shows IPF maps of (a) As-ARB and (b) PA573. The IPF maps constructed from the EBSD measurement data show that fine and ultrafine grains could be observed in As-ARB. Note that EBSD measurement was not performed on PA673 since PA673 seemed to form thick intermetallic layers at the interphase interfaces. This result is reasonable since the conventional ARB process, which is widely used to form ultrafine-grained single-phase metals,24) was adopted in this study. Coarser grains were formed in PA573 compared with As-ARB, as shown in Fig. 6(b). The reason why relatively fine grains remained is that the A5052 layers contained a high concentration of alloying elements, and therefore, grain growth was prohibited by them.
IPF maps of (a) As-ARB and (b) PA573. It is noted that ND and RD are x and y axis, respectively.
In general, the obtained grain size formed by ARB or other severe plastic deformation (SPD) processes is strongly affected by the type and concentration of impurity or alloying elements, in other words, higher concentration of impurity or alloying elements results in smaller grain size. There are other factors which affect to the obtained grain size, such as, initial grainsize before SPD processes, initial orientation of grains, and what element is used as matrix. Thus, it is quite difficult why the grain size of A5052 is finer than that of Cu in As-ARB. For instance, the initial grain size was not controlled to be equal in this study, but, it may affect the obtained grain size of Cu and A5052 in As-ARB. This point is out of focus of this study. However, it should be studied in future if dissimilar metallic laminates are fabricated by the ARB process.
As shown in Fig. 2, the brittle fracture of PA673 was confirmed during the compression test. Figure 7 shows the optical micrograph (a) before and (b) after the drastic drop of flow stress. As can be seen from Fig. 7(a), there is no clear plastic deformation including localized deformation. Clear fracture can be seen in Fig. 7(b), where after the drop of flow stress occurs. The fracture can be categorized into two types; one is delamination of interface formed by the ARB process, and the other is diagonal to CD. The formation of an intermetallic compound layer was also detected in PA573, as shown in Fig. 5. A thicker intermetallic compound layer is expected to have possibly developed in PA673 since the annealing temperature was higher than that for PA573. Therefore, the brittle fracture at the nominal stress at approximately 800 MPa is primarily attributed to the thick intermetallic compound layer formed by annealing at 673 K. In contrast, PA573 did not show brittle fracture, while it had an intermetallic compound layer. As mentioned above, the thickness of the intermetallic layer and original thickness of the As-ARB layers were approximately a few microns and 10 µm, respectively. In such a case, brittle fracture like PA673 does not occur. It can be thought that the localized deformation was formed in PA673 if its deformability is enough, but, thick intermetallic layers did not have enough deformability showing diagonal localized deformation.
Optical micrographs of P673 specimen (a) before and (b) after drastic drop of flow stress is observed during compression test. The white arrows represent the denomination of the interface.
Here, we would like to discuss the growth of the intermetallic layers in terms of interdiffusion of Cu and Al. Based on Table 2, if the interdiffusion constant D of Cu within Al and that of Al within Cu are adopted as 2.5 × 10−15 m2/s and 5.6 × 10−20 m2/s,25) respectively, then the average diffusion distances $\sqrt{Dt} $ for 3.6 ks at 673 K for Cu within Al and Al within Cu are obtained as 3.0 µm and 14 nm, respectively. The estimated $\sqrt{Dt} $ with a maximum of up to ∼5 µm is comparable to that obtained from SEM/EDS analysis. Additionally, the values of $\sqrt{Dt} $ for 3.6 ks at 573 K for Al within Cu and Cu within Al are obtained as 0.35 µm and 0.8 nm, respectively. The estimated $\sqrt{Dt} $ with a maximum of less than 1 µm is again comparable to that obtained from SEM/EDS analysis. Thus, the formation of intermetallic layers is expected to be dominated by the diffusion of Cu in the Al alloy layer in both the 573 K and 673 K annealing. Nevertheless, PA573 showed kink formation and had sufficient plasticity during the compression test since the thickness of the intermetallic compound layer was maintained at less than a few microns after annealing at 573 K. In contrast, PA673 did not have sufficient plasticity during compression due to the micrometer-order-thick intermetallic compound layer.
The formation and propagation of the dark line or region inclined approximately 45° from the compression axis are considered to correspond to development and propagation of an ortho-type kink, based on the in-situ optical microscope images obtained during the compression tests, the analysis using the DIC technique shown in Fig. 4, and previous literature.20) Focusing on the εxy maps in Fig. 4, the absolute value of εxy in the kink band of PA573 was remarkably larger than that of As-ARB. Furthermore, the distribution of εxy was significantly discontinuous and highly concentrated within the kink band in PA573. The difference must be due to the formation of an intermetallic compound layer and grain coarsening induced by annealing at 573 K prior to the compression test. In PA573, annealing at 573 K might improve the strength of the interphasecccccc interfaces, and consequently, a sharp ortho-type kink band was developed.
It can be said that there is a difference in s-s curves between As-ARB and PA573 in terms of work hardening. When Fig. 2 is carefully checked, As-ARB shows almost no work hardening after kink deformation, but PA573 seems to show slight work hardening. Considering the data seen in Fig. 4, there is a difference between As-ARB and PA573, such as, fraction of localized εyy region is larger in PA573 than that in As-ARB. It could be attributed to the slight work hardening seen in s-s curve of PA573.
Furthermore, there might be a small difference between As-ARB and PA573 before kink deformation occurs, since the clear microstructural difference, such as, the presence of intermetallic layers at interphase-interface in PA573, exists. When the s-s curves shown in Fig. 2 is carefully checked, the gradient (apparent Young’s modulus) is different among three types of specimens. If specimens are arranged in order of apparent Young’s modulus from largest to smallest one, it becomes P673, P573 and As-ARB. Thus, the presence of thin intermetallic layers at interphase-interphase may affect the apparent Young’s modulus of the DMLs before kink deformation occur.
In the case of LPSO-Mg alloys, the following empirical criteria for obtaining kink strengthening are satisfied: (1) the material consists of an alternating structure with hard and soft layers in terms of the shear modulus, (2) the thickness of each layer is on the order of submicrometers, (3) the active slip systems are restricted, and (4) delamination between the hard and soft layers does not occur during kink formation. Hereafter, whether the criteria are valid for the kink formation in the Cu/A5052 DMLs in the present study is discussed.
The Cu/A5052 DMLs satisfy the first criterion since the MFS consists of high (Cu) and low (A5052) shear modulus layers. However, the second criterion is not satisfied since the thickness of each layer is on the order of micrometers. The third criterion is also not satisfied in the present study since the Cu and A5052 alloys have a face-centered cubic structure and the slip system is not restricted in the plane parallel to the layers. The fourth condition is apparently satisfied for As-ARB and PA573 since a rigid interphase interface is required; otherwise, the layers are peeled off during the compression tests. The criteria are discussed in more detail below.
Here, the first and second criteria are reconsidered. Note that strata having extremely large thicknesses on the order of kilometers also show kink band formation.19) Recently, Hagihara et al. reported this kind of deformation for Al/Al2Cu eutectic alloys having layers with thicknesses larger than 1 µm.22) Therefore, the thickness ratio of the hard and soft layers might be more important than the actual thickness of each layer.26,27)
Next, the third condition is reconsidered. The Cu/A5052 DMLs were fabricated by the ARB process in this study, and ultrafine grains were formed via “grain subdivision28)”, which is also observed in the ARB-processed single-phase material.29) In general, ARB-processed materials do not show high work hardening ability since an ultrafine-grained microstructure is formed via grain subdivision. Thus, the usual uniform deformation due to the introduction of dislocations via plastic deformation is suppressed. This is probably why kink deformation appeared in the Cu/A5052 DMLs. In other words, uniform deformation via dislocation slip-mediated plasticity must be prohibited regardless of the reason when kink deformation is required.
PA573 also shows kink deformation since the work hardening ability of each layer may still be suppressed. Furthermore, the intermetallic compound layer is not thick enough to allow brittle fracture. Therefore, the criteria for the ortho kink and kink bands to appear in a DML with the MFS can be rewritten as follows.
The preannealing process changes the morphology of the kink band, as shown in Figs. 3 and 4. The kink boundaries represented by red dashed lines seem much sharper for PA673 (Fig. 3(c)) than those for PA573 (Fig. 3(b)). Figure 4 also shows a sharp jump of εxy around the kink boundaries in PA573 compared with the As-ARB sample. The difference between the As-ARB and PA573 samples is whether preannealing was applied before the compression test. When we remember the rewritten criteria for a DML with the MFS, the difference in the morphology of the kink band can be understood. The preannealing at 573 K forms a thin intermetallic compound layer at the interphase interface. Obviously, the intermetallic compound layer is brittle, and therefore, PA573 cannot uniformly deform compared with the As-ARB sample. The εyy map in Fig. 4 shows that the As-ARB sample shows a relatively spatially homogeneous higher εyy region outside of the higher εxy region. In contrast, PA573 shows a spatially inhomogeneous higher εyy region outside of the higher εxy region. It should be pointed out the there is a possibility that such strain inhomogeneity shown in Fig. 4(i) could be attributed to either the sample geometry or compatibility on both sides across the kink boundary.
Since shaper kink boundary is formed in P573 compared with As-ARB, homogeneous deformation along the vertical axis (εyy) is requested around ortho-kink in terms of compatibility on both sides across the kink boundary. It could be the reason why slight work hardening can be seen on s-s curve of P573 shown in Fig. 2. There is a possibility that P573 is better to be treated as DML with the MFS composed by more than two types of layers, Cu, A5052 and metallic compounds layers. Thus, further investigation on DML with the MFS is requested.
Cu/A5052 DMLs with ∼10 µm thick layers were fabricated by the ARB process. The composites with and without preannealing were compressed, and some composites showed ortho kink and kink band formation. The following are the criteria for kink formation in DMLs.
This work was supported by the Japan Society for the Promotion of Science (JSPS) KAKENHI for Scientific Research on Innovative Areas “MFS Materials Science” (Grant Number: JP18H05481, 19H05123, and 21H00097). The authors would like to thank Prof. C. Watanabe in Kanazawa University for SEM observations and EBSD measurements.
Author contributions
MY: writing the original draft. MY and YM: performing mechanical tests and microstructure observations and preparing all figures. IK and YM: conceptualization, visualization and project administration. IK, TF, and YM: writing – review & editing, supervision, and funding acquisition. All authors discussed the results and commented on the manuscript.
