2019 Volume 60 Issue 2 Pages 254-262
The mechanical properties and nanostructure of the multi-layered aluminum alloy sheet were investigated by tensile test, hardness test, electron probe microanalysis and micro-small-angle X-ray scattering in scanning mode, focusing on the distributions through the thickness. The multi-layered sheets consisting of highly concentrated Al–Mg and Al–Zn alloys show a remarkable increase in proof stress after interdiffusion and artificial aging. The predominant layers to contribute to the proof stress change from the layers with higher Zn/Mg ratio in T4 temper to the layers with lower Zn/Mg ratio after artificial aging. These age-hardening responses depend on the layers, which are large in the layers with higher Zn/Mg ratio, whereas small in the layer with lower Zn/Mg ratio. These noticeable bulk properties are ascribed to the local change in the types, volumes and morphologies of the G.P. zones and/or metastable phases depending on the concentration profiles through the thickness, which are produced from these unique multi-layered structures.
Aluminum alloys are widely used to variety of structural applications for their good balance of moderate strength, satisfactory ductility and corrosion resistance as well as high specific strength. Especially in the automotive industry, the increasing demand for the reduction in the emissions of carbon dioxide has promoted the application of aluminum alloys to the steel body structure in order to reduce the vehicle body weight. One of the technical issues is the improvement of balance in these properties at a higher level, because aluminum alloys are inferior to high tensile steels in these balances. In particular, the increase in the strength brings about the decrease in the ductility and corrosion resistance for the conventional aluminum alloys.1–3) One of the reasons for the difficulty to overcome these subjects is the limitation of the homogeneous microstructure with an inherent lack of diversity.
Recently, in order to break through these limitations, a new approach of material design strategy of ‘architectured materials’ is introduced.4,5) In this approach, in order to obtain materials with unusual combinations of properties, heterogeneous microstructural design and optimization of geometry is combined. In aluminum alloys, one of the typical applications for this approach is layered products,6) such as aluminum heat exchangers for vehicles.7) In these heat exchangers, side materials for brazing or sacrificial protection are roll bonded onto a core material for forming, which enables the layered sheet to play various functions. Another application of the layered sheet is automotive sheet for closure panels. The layered sheets for good bendability, weldability or corrosion resistance are produced for several high performance applications.6)
The concept of heterogeneous microstructural design for these layered sheets is categorized in two groups. One is the concept that these layered sheets are designed to reduce mutual diffusion across the layers in their production process in order to maintain their initial compositions and prevent deterioration of their original efficiency. Most of the layered products described above belong to this concept. Another is that the mutual diffusion is promoted to take advantage of the interaction of the constituent elements among the layers. The latter concept can be an effective way of increasing strength through a combination of interdiffusion of the selected constituents and their precipitation hardening to maximize their effects. In some types of heat exchangers, age hardening is utilized by the diffusion of Mg from the sacrifice protection layer to the core layer and subsequent interaction with Si in the core layer, resulting in the formation of Mg–Si precipitates.7) The extent of strengthening in the heat exchangers by this approach is, however, small because the solubility limit of Mg2Si phase in Al–Mg–Si ternary system is relatively small. Whatever the case maybe, this approach has not been sufficiently studied yet. Thus, by selecting appropriate alloy system and making the most of its age hardening potentials, it can be possible to improve the balance between strength and other mechanical and chemical properties at a higher level.
In this study, we aim to conduct a basic research on the strengthening of the multi-layered sheet by the unique configuration of the layers and interdiffusion of the constituent elements between them. We select highly concentrated Al–Mg and Al–Zn binary alloys as basic alloy systems which compose the multi-layered sheets. These alloys are cladded alternately and annealed to interdiffuse, resulting in the formation of Al–Zn–Mg ternary alloy system. As for this alloy system, G.P. zones and several types of metastable/stable precipitates (G.P. zones, η′, η, T′ and T phases) form during aging process, which contributes to the age hardening.8) Many attempts have been made to characterize these precipitates by transmission electron microscopy (TEM), atom probe tomography (APT) and small angle X-ray scattering (SAXS) to relate the microstructural factors to the mechanical properties.9–18) Most of these previous studies were, however, focused on the conventional Al–Zn–Mg alloys where both of Mg and Zn additions are smaller than about 3 at% and the zinc-to-magnesium concentration ratio (Zn/Mg ratio) is smaller than about 2.9–18) This paper, on the other hand, describes the results of the multi-layered alloys which have higher concentration of Mg and Zn than the conventional alloys. Due to the wide range of the compositional gradient through the thickness after diffusion annealing, the multi-layered sheets form a variety of precipitates through the thickness, which brings about the functionally graded properties and nanostructure at a macroscopic scale in the multi-layered sheets. To evaluate these characteristically heterogeneous distributions of the precipitates through the thickness, micro-small-angle X-ray scattering (SAXS) in scanning mode is applied.
An aluminum alloy with the five-layered structure, where two types of aluminum alloys are alternatively accumulated, was studied. Table 1 shows the main chemical compositions of these layers with minor additions of Zr to control grain structure. As shown in Fig. 1, the first, third and fifth layers stacked from the surface correspond to Al–5.5Mg, whereas the second and fourth layers correspond to Al–15Zn–0.5Cu. All these layers are of uniform thickness. The production process is described below. The ingots of Al–5.5Mg and Al–15Zn–0.5Cu were homogenized at 723 K for 28.8 ks and hot-rolled to the same thickness. The surface of these hot strips was polished by wire brushes to remove the surface oxide film before roll bonding. After stacking of these hot strips as shown in Fig. 1, the multi-layered plate was bonded by rolling at 573 K. Subsequent cold rolling to the final thickness of 1 mm and annealing at 673 K for 7.2 ks were performed for the plate (O temper). Heat treatment at 743 K for 24 ks was conducted for the multi-layered sheets to cause both interdiffusion of the constituent elements across these layers and solid solution of them. After these heat treatments, the sheets were quenched into water and naturally aged at room temperature for 1.2 Ms (T4 temper), which was followed by artificial aging at 398 K for up to 86.4 ks.
Configurations of the aluminum alloy sheet with five-layered structure.
Mechanical properties of these sheets were evaluated by tensile test with the initial strain rate of 3.33 * 10−3 s−1. As the multi-layered sheets have inhomogeneous distributions of the constituent elements and nanostructure through the thickness, characterization of these distributions was performed by several measurements and analyses. The concentration profile measurement was made by electron probe microanalysis (EPMA), hardness distribution was examined by micro Vickers hardness test. The EPMA profiles for each element were measured at intervals of 1 µm through the thickness and the average concentrations were calculated at intervals of 10 µm by using all the four data up and down the corresponding position which was set at intervals of 10 µm through the thickness. The values of the micro Vickers hardness were measured at intervals of 20 µm through the thickness. In this study, scanning micro-SAXS was performed to characterize the nanoscale precipitates in the multi-layered sheets at BL40XU of SPring-8. The X-ray, which has incident energy of 15 keV, is shaped by a pinhole with a diameter of 5 µm. The intensity distributions of SAXS were measured through the whole thickness of the multi-layered sheets at intervals of 10 µm in T4 temper condition and 25 µm after artificial aging conditions, respectively. To check the exact positions of the measurements in the sample, the intensity of Zn–K fluorescent radiation was also monitored with a silicon drift detector (SDD) simultaneously.
From the measured spectra, the gyration radius, Rg was determined by Guinier approximation.19) The distribution of the integrated intensity, Q was calculated from eq. (1).20)
| \begin{equation} Q = \int q^{2}I(q)dq = \Delta\rho^{2}V_{f}(1 - V_{f}) \end{equation} | (1) |
The interparticle distance between nearest precipitates, Lp was also estimated from the spectra. From Rg and Lp, Vf was estimated according to eq. (2).20)
| \begin{equation} V_{f} \approx (4\sqrt{2\pi}/3)(R_{g}/L_{p})^{3} \end{equation} | (2) |
Figure 2 shows the concentration distribution of the multi-layered sheets after the annealing at 673 K for 7.2 ks (a), and the diffusion annealing at 743 K for 24 ks (b). The concentration profile of each constituent element gives an interdiffusion region. The length of the interdiffusion region d(X) is obtained from the slope of the tangent line at half of the concentration of the constituent element X to the original layers, as shown in Fig. 2(a). From these analyses, d(Zn), the length of the interdiffusion region of Zn content, at the both interfaces of the second layer in the multi-layered sheet in O temper is estimated at 115 µm. This result indicates that this annealing condition does not promote the interdiffusion of these constituent elements in this multi-layered sheet. As for the multi-layered sheet in T4 temper, the outer and inner interfaces of the second layer give interdiffusion regions with 677 and 903 µm length, respectively. These results indicate that the constituent elements of these layers almost diffuse to the both surfaces in T4 temper. The difference in the interdiffusion length between the both interfaces of the second layer is supposed to be caused by the asymmetric configuration of the second layer; outer interface (surface side) of the second layer comes in contact with only the first layer, whereas inner interface (center side) of the second layer is faced to the third layer which is also adjacent to the forth layer. Thus the inner interface of the second layer is presumed to be influenced by the interdiffusion of Zn element from the fourth layer through the third layer for the long time annealing condition, resulting in the increase in the apparent interdiffusion length of the inner interface of the second layer. Consequently, these characteristic concentration profiles through the thickness affect to the hardness distributions in addition to the precipitation distributions. The concentration of Cu, on the other hand, varies from 0.13 to 0.37 at% through the sheet thickness after diffusion annealing at 743 K for 24 ks. The variation is much smaller than those of other constituent elements.
Concentration distributions of Mg, Zn and Cu elements and concentration ratios of Zn to Mg through the thickness of the present multi-layered sheets, (a) after annealing at 673 K for 7.2 ks (O temper), (b) after diffusion annealing at 743 K for 24 ks (T4 temper).
The hardness distributions through the thickness of the multi-layered sheets, which are heat-treated in T4 temper and artificially aged at 393 K, is shown in Fig. 3. As for T4 temper, there are two regions of these hardness distributions through the thickness. The hardness, which shows the minimum value at the surface, drastically increases to its peak value at the position of about 250 µm from the surface, then, slightly decreases at the inside positions, which continues to the center of the thickness. Furthermore, the artificially aged sheets show the complex changes. With increase in the artificial aging time, the hardness increases in almost all the positions through the thickness. However, the artificial age hardening responses depend on the positions. In the position range of 200–400 µm from the surface, a slight increase in the hardness during the artificial aging is observed, whereas especially in the position ranges of about 100–180 µm and 400–500 µm, significant increases occur, which results in the shift of the maximum peak to the surface side. To clarify quantitatively the dominant layers in the age hardening, the multilayered sheet was divided into ten segments and the ratio of the integrated value of the Vickers hardness in each layer to the sum of all the Vickers hardness throughout the thickness was calculated, respectively.
Hardness distributions through the thickness of the multi-layered sheets heat-treated in the T4 temper and artificially aged at 393 K.
Figure 4 shows the proof stress during artificial aging obtained by the tensile test. In this figure, the stacked column chart indicates the contributions of these ten layers to the proof stress, as mentioned above. The proof stress increases significantly during the artificial aging and reaches about 580 MPa after aging at 393 K for 86.4 ks. The contributions of these layers to the artificial age hardening are roughly divided into the following three groups. The first group of the layer C and D with higher Zn/Mg ratio, which has a major contribution to the proof stress in T4 temper, decreases in the proportion to the proof stress as the artificial aging proceeds. On the other hand, the second group of the layer B and E with lower Zn/Mg ratio increases in the proportion to the proof stress during artificial aging, resulting in a predominant contribution to the increase in the proof stress. The third group of the layer A with the minimum Zn concentration has the smallest proportion to the proof stress and keeps almost constant during artificial aging. These results indicate that the macroscopic concentration profiles of the constituent elements through the thickness intricately affect to the macroscopic profiles of the mechanical properties during artificial aging, which is attributed to the inhomogeneous distributions of G.P. zones and/or metastable precipitates through the thickness.
The proof stress during artificial aging obtained by the tensile test. The stacked column chart indicates the contributions of the ten layers to the proof stress (The stacked columns are composed of the five segments because the two equivalent layers, which are placed symmetrically to the center of the thickness of the multi-layered sheets, are merged respectively).
The distributions of the radius of gyration Rg and integrated intensity Q through the thickness of the multi-layered sheets in T4 temper and after artificial aging at 393 K for 86.4 ks are shown in Fig. 5. G.P. zones and/or metastable precipitates have characteristic distributions through the thickness of the multi-layered sheet. Both the Rg and Q in T4 temper have maximum values between the layer C and D, and minimum values in the layer A, which corresponds to the Zn/Mg ratio. All the Rg through the thickness increase during artificial aging. All the Q, whereas, decrease after artificial aging. The amount of decrease in Q is remarkable, especially in the layer C and D, which have higher values in Q than other layers. As Q is affected by the electron density contrast Δρ and volume fraction Vf according to the eq. (1), it is considered that these changes in Q would reflect the changes in the type and amount of the metastable phases. The Vf in T4 temper and after artificial aging is shown in Fig. 6. The data of the positions of less than 100 µm after artificial aging condition were eliminated because these Vf could not be evaluated accurately due to their intensity profiles of SAXS. The distributions of the Vf show the distinguishing profiles through the thickness. In T4 temper, the layer B and E with lower Zn/Mg ratios have two large peaks of the Vf. These distributions of the Vf through the thickness, however, don’t explain the hardness distributions with higher values at the layer C and D shown in Fig. 3, and the higher proportion to the proof stress of the layer C and D shown in Fig. 4 in T4 temper well enough. The Rg distributions in T4 temper would rather correspond to these distributions of hardness and mechanical property. After the artificial aging at 393 K for 86.4 ks, all the Vf through the thickness decrease. There also remain two peaks of the Vf in the layer B and E, which have smaller heights compared to those in T4 temper.
Distributions of the radius of gyration Rg and integrated intensity Q obtained by scanning SAXS measurements.
Distributions of volume fraction Vf obtained by scanning SAXS measurements.
From those results, these profound effects on the mechanical strength of the multi-layered sheets in the present study are due to the characteristic variations of the size, volume fraction of the second phase particles such as G.P. zones and/or metastable precipitates through the thickness directions and during the aging process.
As described in the section 3.1 and 3.2, each layer shows the various types of age hardening for the multi-layered sheets. It is considered that these differences are attributed to the inhomogeneous distributions of G.P. zones and/or metastable precipitates through the thickness, which corresponds to the wide range of the matrix compositions after interdiffusion of the constituent elements during the diffusion annealing as noted in the section 3.1 and 3.3. Thus, the understanding of the precipitate phases on the equilibrium and non-equilibrium phase diagrams gives some information about these age hardening behaviors although there are few reports about non-equilibrium phase diagrams of Al–Zn–Mg ternary systems. In this section, the discussion about the precipitate phases is based primarily on the equilibrium phase diagrams. Figure 7 shows the isothermal section of Al–Zn–Mg ternary phase diagram at 393 K calculated by Thermo-calc software with Al alloy database version TCAl8. The representative compositions which correspond to the positions from the surface to 500 µm in the thickness direction of the multi-layered sheets after the diffusion annealing at 743 K for 24 ks are displayed as open and closed circles in this diagram. It is notable that the compositions of this multi-layered sheet spread across the two dual-phase and two triple-phase coexisting regions, which corresponds to the wide range of the Zn/Mg ratio. As for the decomposition sequences of the aluminum supersaturated solid solution, it is known to form mainly metastable η′ phase in the region of moderate Zn/Mg ratio (between about 1 and 2) and metastable T′ phase in the region of lower Zn/Mg ratio (less than about 1) after the formation of G.P. zones in the Al–Zn–Mg alloys.9,18,20) As for this multi-layered sheet, furthermore in the region of higher Zn/Mg ratio where α, η and Zn phases coexist in the equilibrium phase diagram, metastable α′R phase, which is usually formed in the Al–Zn binary system,21) would be formed. As shown in Fig. 7, it is clear that almost all the compositions through the thickness are put on the same straight line represented as the equation; X(Zn) + 3.5X(Mg) = 14, where X(Zn) and X(Mg) indicate the mole percent of Zn and Mg, respectively. Thus, to estimate the thermodynamic stability of these phases, the vertical section of this equilibrium phase diagram is calculated according to this equation, as shown in Fig. 8. The solvus temperatures for G.P. zones20,22,23) plotted in this phase diagram are above room temperature in all the range of Zn/Mg ratio in this figure, which suggests that at least G.P. zones are formed in all the layers of the multi-layered sheet in T4 temper. As shown in Fig. 5, the inhomogeneous distribution of the integrated intensity Q in T4 temper, which is a function of Δρ and Vf, corresponds to the variation of the type and composition in the precipitates or volume fraction through the thickness.
Isothermal section of Al–Zn–Mg ternary phase diagram at 393 K calculated by Thermo-calc. The open and closed circles indicate the representative compositions of the positions at 50–300 µm and 300–500 µm at 50 µm intervals in the thickness direction of the multilayered sheets after diffusion annealing at 743 K for 24 ks, respectively. The broken line and chain line indicate the calculated solubility limits of aluminum primary solid solution equilibrating with α2 and GP zone, respectively at 383 K.20)
Vertical section of Al–Zn–Mg ternary equilibrium phase diagram calculated by Thermocalc. The horizontal axis corresponds to the relation; X(Zn) + 3.5X(Mg) = 14, where X(Zn) and X(Mg) indicate the mole percent of Zn and Mg, respectively. The open circles indicate the compositions of the positions at 50–300 µm and 300–500 µm in the thickness direction of the multilayered sheets after diffusion annealing at 743 K for 24 ks. The closed diamond23) and square20) indicate the calculated solubility limits of aluminum primary solid solution equilibrating with G.P. zone at R.T. and 383 K in Al–Zn–Mg ternary system, respectively. The closed triangle22) indicates the experimental solvus temperature for G.P. zone in Al-14 mole percent Zn alloy.
Figure 9 shows the variation in Δρ with Zn/Mg ratio in T4 temper. The Δρ shows a characteristic trend; it increases with increase in Zn/Mg ratio. The two hypotheses are formed about the reason for this trend. One is that there are two types of G.P. zones which have the structures close to T′ or η′ phase depending on Zn/Mg ratio. The other is that the compositions of G.P. zones vary continuously as a function of Zn/Mg ratio. As for the first hypothesis, analysis of the composition of metastable precipitates by APT revealed that T′ phase has Zn/Mg ratio of about 0.7,18) whereas η′ phase has the higher Zn/Mg ratio which is slightly larger than 1.10,12,14) On the other hand, Adachi et al. investigated the relation between Δρ and metastable phases in Al–Zn–Mg ternary system by SAXS. They reported that Δρ for T′ phase has variations with Zn/Mg ratio, whereas that for η′ phase has a constant value to Zn/Mg ratio.20) Thus, there is a discrepancy between this hypothesis and Adachi’s result. As for the second hypothesis, the compositions of the G.P. zones and precipitates are investigated by APT. Zn/Mg ratio in the G.P. zones increase in the increase in Zn/Mg ratio of the alloy; the alloys with Zn/Mg ratio of about 1 form G.P. zones with Zn/Mg ratio ranging from 1 to 1.4,9,14) whereas the alloys with Zn/Mg ratio of about 1.9 form G.P. zones with Zn/Mg ratio ranging from 1.57 to 2.16) The Δρ in T4 temper indicates the same tendency as these reports. From these discussions, it is supposed that these trends for Δρ correspond to the compositions of G.P. zones.
Variation in Δρ with Zn/Mg ratio after natural aging (T4).
As shown in Fig. 5 and 6, all the integrated intensity Q and the volume fraction of the precipitates Vf decrease after artificial aging. The reversion of the G.P. zones, which formed in T4 temper, must occur all through the thickness during the artificial aging. The solvus temperature for G.P. zones, as shown in Fig. 8, is located below the artificial aging temperature in the region of lower Zn/Mg ratio and above the artificial aging temperature in the region of higher Zn/Mg ratio. Thus, it is presumed that full and partial reversions take place in the region of lower and higher Zn/Mg ratio, respectively. The drastic decrease in Q in the layer C and D with higher Zn/Mg ratio would be attributed to the reversion of the metastable α′R phase, which is formed in the region where α, η and Zn phases coexist in the equilibrium phase diagram.
After artificial aging, metastable T′ and η′ phases could form in the multi-layered sheets with wide variety of compositions. To characterize these metastable phases in detail, another information such as TEM is needed in addition to SAXS, which is an issue in the future.
4.2 Effect of the Zn/Mg ratio on the strength mechanismAs noted in the section 3.2 and 3.3, G.P. zones or/and metastable phases with several sizes and volume fractions have influence on the strength for the multi-layered sheets. In this section, the strength mechanism is discussed below.
In order to analyze the contribution of the G.P. zones and meta-stable phases to the natural and artificial age hardening quantitatively, the amount of the solid solution strengthening after the heat treatment for interdiffusion should be evaluated because these distinctive hardness distributions are the results of the combinations of the solid solution hardening and the precipitate hardening. As the actual measurement of the hardness distribution for solid solution after the heat treatment for interdiffusion is quite difficult due to its natural age hardening, the calculated hardness for solid solution is derived by using the following two equations. It is reported that solid solution hardening is dependent on the square root of the solute concentration (at%) C, namely Δσ ∝ A$\sqrt{C} $.24) The values of the coefficient A are calculated by the first principle calculation;25) the coefficients for Mg, Zn and Cu elements are 25.2 MPa/$\sqrt{C} $, 6.3 MPa/$\sqrt{C} $ and 51.9 MPa/$\sqrt{C} $, respectively. It is reported that the relationship between hardness and proof stress in aluminum alloy is expressed by eq. (3).26)
| \begin{equation} \mathit{HV} = 13.4 + 7.69\sqrt{\sigma - 36} \end{equation} | (3) |
Distributions of the hardness corresponding to solid solution (SS), and the hardness excluding the effect of solid solution hardening after natural aging (T4) and after artificial aging at 393 K for 86.4 ks (86.4 ks).
The amount of age-hardening from solid solution ΔHV at each position as a function of the microstructural parameters based on (a) cut-through mechanism and (b) by-pass mechanism.
Using the amount of age-hardening from solid solution (ΔHV) estimated as above, two strength models are evaluated. According to Gerold and Harberkorn’s coherency strain model, the cut-through mechanism is represented by eq. (4).28)
| \begin{equation} \tau\propto \mu |\varepsilon|^{3/2}\sqrt{\frac{V_{f}\cdot r}{b}} \end{equation} | (4) |
| \begin{equation} \tau\propto \mu b\frac{\sqrt{V_{f}}}{r} \end{equation} | (5) |
In Fig. 11(a), one characteristic aspect is recognized about ΔHV; ΔHV after artificial aging is larger than that in T4 temper in the range of the same radius and volume fraction of the precipitates. This means that the precipitates after artificial aging become harder than that in T4 temper by the change in the interaction of precipitates with dislocation. As shown in Fig. 6, the volume fraction of the precipitates decreases all through the thickness after the artificial aging. This negative effect of the decrease in the volume fraction of the precipitates must be canceled out and exceeded by the strengthening effect.
Furthermore, there are two more remarkable features in Fig. 11(a) about the slope of each group of data which corresponds to the coherency strain. The first point is the difference between the layers; the coherency strain at the layer B and E is larger than that at the layer C and D. The coherency strain is defined as the difference of average atomic size between the precipitate and the matrix.30) The atomic mismatch between Mg and Al is larger than that between Zn and Al,31) which causes the increase in the coherent strain of the G.P. zone with lower Zn/Mg ratio. As for the metastable phases, T′ phase with higher Mg concentration has larger coherent strain than η′ phase with lower Mg concentration does.32) These composition dependences on the atomic mismatch between the matrix and clusters or precipitates would bring about the higher coherency strain in the layer B and E with lower Zn/Mg ratio. The second point is the difference between aging conditions; the coherency strain of the precipitates in T4 temper decreases after artificial aging. This could be attributed by the change of the types of precipitates. As described in the section 4.1, G.P. zone is dominant in T4 temper, whereas the η′ and/or α′R phase is formed and coexists with G.P. zone after the artificial aging. G.P. zone has coherent interface between the matrix and itself, while the η′ and/or α′R phase has semi-coherent interface, resulting in the decrease in the coherency strain. These changes of coherency strain would bring about the low dependency of hardness on the microstructural parameters after the artificial aging.
In the present study, it reveals that the multi-layered sheets with compositional gradient through the thickness show profound effects on the mechanical properties by controlling the interdiffusion and age hardening behavior of the layers which initially have different compositions. These properties can be obtained by the optimization of the distribution of G.P. zones and/or precipitates, which have different sizes and volume fractions through the thickness. The multi-layered sheets also have characteristic distributions through the thickness, such as composition and hardness. These distributions can potentially affect to the chemical and mechanical properties to improve the corrosion resistance and ductility, which is the main subject of the high strength aluminum alloys. Further studies are needed to achieve these properties and in progress.
The mechanical properties and nanostructure of the multi-layered aluminum alloy sheet were investigated by tensile test, hardness test, EPMA and scanning micro-SAXS, focusing on the distributions through the thickness. The obtained results are as follows.
This research is based on results obtained from a project commissioned by the New Energy and Industrial Technology Development Organization (NEDO). The synchrotron radiation experiments were performed at the BL40XU of SPring-8 with the approval of the Japan Synchrotron Radiation Research Institute (JASRI) (Proposal No. 2014B1597, 2015A1684, 2015B1597, 2016B1580 and 2017A1597).
