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Al clad sheets through annealing and secondary rolling process
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Improvement of interfacial bonding strength in roll-bonded Mg/Al clad Sheets through annealing and secondary rolling process Jung-Su Kim, Kwang Seok Lee, Yong Nam Kwon, Byeong-Joo Lee, Young Won Chang, Sunghak Lee
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S0921-5093(15)00051-9 http://dx.doi.org/10.1016/j.msea.2015.01.035 MSA31973
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Received date: 30 October 2014 Revised date: 13 January 2015 Accepted date: 14 January 2015 Cite this article as: Jung-Su Kim, Kwang Seok Lee, Yong Nam Kwon, ByeongJoo Lee, Young Won Chang, Sunghak Lee, Improvement of interfacial bonding strength in roll-bonded Mg/Al clad Sheets through annealing and secondary rolling process, Materials Science & Engineering A, http://dx.doi.org/10.1016/j. msea.2015.01.035 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Improvement of Interfacial Bonding Strength in Roll-Bonded Mg/Al Clad Sheets through Annealing and Secondary Rolling Process Jung-Su Kima, Kwang Seok Leeb, Yong Nam Kwonb, Byeong-Joo Leea, Young Won Changc, and Sunghak Leea,* a
Center for Advanced Aerospace Materials
Pohang University of Science and Technology, Pohang, 790-784, Korea b
Materials Deformation Department, Light Metal Division
Korea Institute of Materials Science, Changwon, 641-010, Korea c
Graduate Institute of Ferrous Technology
Pohang University of Science and Technology, Pohang 790-784, Korea +82-54-279-2140(Telephone) / +82-54-279-5887(Fax) *Corresponding Author e-mail: [email protected] (S. Lee)
Abstract The effect of annealing and secondary rolling process has been investigated in relation to evolution of interface microstructure as well as interfacial bonding mechanisms of roll-bonded two-ply Mg/Al clad sheets in this study. Two types of thin reaction layers, viz., γ (Mg17Al12) and β (Mg2Al3) phase layers, were observed to form along the Mg/Al interface, the thickness of which was found to increase with annealing time. The grains in the γ layer were grown coarsely in a columnar shape, as they grew continuously in coarse-grained Mg substrate alloy side. The thickness of the γ layer was found to reduce significantly after the secondary one-pass rolling. At the same time, the initially columnar γ grains became refined into equi-axed grains by recrystallization caused by large plastic deformation during the secondary rolling. The serrated flow observed in the reaction layers during a nano-indentation test in the γ layer disappeared completely without any sign of micro-cracks after the secondary rolling. This in turn implies that the γ layer, acted as an embrittlement site before the rolling, may not be brittle any longer after the rolling. The grain refinement and disappearance of columnar grains in the γ layer appear to greatly improve the interfacial bonding strength in these clad sheets. This secondary rolling process seems to provide a great advantage in terms of manufacturing, since the same rolling stand for roll-bonding process can also be used without any additional equipment. Keywords: Mg/Al clad sheet, annealing, secondary one-pass rolling, interfacial bonding, γ (Mg17Al12) phase, β (Mg2Al3) phase
1. Introduction Multi-layered metallic clad sheets combining two or more alloys have recently been developed to provide a new multi-functional property utilizing characteristic properties of each metallic alloy. Disadvantages of one alloy sheet can therefore be effectively complemented or improved to provide entirely new application areas in various industries such as automotive, aerospace, and kitchen utensils industries, and so on [1-5]. These multi-layered clad sheets have generally been fabricated by using Al alloys, Ti alloys, or stainless steels to utilize their excellent heat conductivity and corrosion resistance as well as other specific mechanical properties [6, 7]. Global warming has especially become a serious issue all over the world including carbon dioxide emissions to generate a significant attention in structural materials field. In order to reduce carbon dioxide emission, a number of research and development efforts have actively been conducted on lightweight materials such as Mg and Al alloys [8-10]. Mg alloys are the lightest-weight metal commercially available today, having various advantageous properties in terms of specific strength, damping capacity, electromagnetic shielding, and dimensional stability [11, 12]. Actual industrial application areas of Mg alloys are, however, quite limited due to their inherent drawbacks in relation to poor corrosion resistance, formability, and surface quality. In order to expand the applications of Mg alloys, various types of clad Mg alloys have been developed by bonding with other alloys such as stainless steels and Al alloys to complement the drawbacks of Mg alloys [13-16]. Direct solid-state bonding between Mg and Al alloys is, however, seriously restricted due to limited formability of Mg alloys at low and intermediate temperatures [15]. Various solid-state bonding processes of Mg/Al sheets have been developed to date including roll-bonding, friction-bonding, diffusion-bonding, and explosion-bonding [17-23]. Among these, roll-bonding is the most widely used process for fabricating Mg/Al clad sheets, considering economic advantages in terms of convenient utilization of conventional rolling facilities, productivity, and easy combination of clad metals [24]. A warm roll-bonding is generally used to circumvent the poor room-temperature formability of Mg alloys followed by subsequently annealing treatment to control mechanical properties of clad sheets. Brittle reaction phases are, however, inevitably formed along the solid-state-bonded interface in the form of intermetallic compounds, viz., γ (Mg17Al12) and β (Mg2Al3) phases, as readily expected from the Mg-Al binary phase diagram [25, 26], to deteriorate overall mechanical properties as well as interfacial bonding strength. These brittle reaction layers have recently been reported by Lee et al. [27] to act as interfacial delamination sites during the mid-stage flexural bending test of a three-ply laminated composite of stainless steel/Al/Mg. A thin interlayer of pure Zn was inserted between Mg and Al alloy sheets to prevent this interfacial delamination problem as suggested by Zhao et
al. [28] The insertion of a high-purity Zn foil was found to retard or suppress the formation of γ and β phases during the solid-state diffusion bonding. This method may not be suitable for the roll-bonding process of wide-width clad sheets in view of manufacturing cost. The size, morphology, and thickness of reaction phases should be optimally controlled in order to improve the interfacial bonding strength of Mg/Al clad sheets. The effects of these parameters on interfacial bonding strength have not, however, been clarified to date, except the generally accepted fact that interfacial bonding strength could be enhanced by controlling the thickness of reaction phases to a certain level [29-32]. Elucidation of mechanisms in relation to microstructural modification is required as well, since the only variation of thickness of reaction layer cannot provide a plausible explanation on interfacial bonding strength [16]. In addition, a cost effective fabrication method of Mg/Al clad sheets should be developed by utilizing conventional facilities such as rolling stands to obtain adequate microstructures and required properties simultaneously. This fabrication idea is quite useful especially in relation to the production costs of Mg/Al clad sheets having excellent microstructures and properties, taking full advantages of individual Mg and Al alloys. Mg/Al clad sheets were therefore fabricated in this study by a roll-bonding process together with a subsequent annealing treatment followed by a warm rolling. This process route provided the merits of economical fabrication of Mg/Al clad sheets and also the improvement of their interfacial bonding strength simultaneously. Interfacial microstructures of fabricated sheets were then analyzed, and their properties were evaluated by conducting nano-indentation as well as roller-drum peel tests. Interfacial bonding mechanisms were also investigated in relation to the improving interfacial bonding strength, by focusing specifically on how the size, morphology, and thickness of reaction phases affected the initiation and propagation of cracks in the Mg/Al interface.
2. Experimental procedures Multi-pass warm-rolling was performed first using sheets of 3004 Al and AZ31 Mg alloys at 400 oC with the reduction ratio of 60 % to fabricate 3.3-mm-thick two-ply clad sheets. The 3004 Al alloy sheet had the thickness of 1.0 mm with the composition of Al-(1.0~1.5)Mn(0.8~1.3)Mg-(≤0.7)Fe-(≤0.3)Si, while the AZ31 Mg alloy sheet had the thickness of 2.3 mm with the composition of Mg-(2.5~3.5)Mn-(0.6~1.4)Zn-(0.2~1.0Mn)-(≤0.1)Si. These roll-bonded clad sheets were then subsequently annealed at 300 oC for 10 min and 60 min before air-cooling. Right after this annealing treatment, one-pass warm-rolling (300 oC) was carried out for some of the sheets with the total reduction ratio of 22 % or 41 %. The as-roll-bonded clad sheet is designated as ‘AR’, while the sheets annealed at 300 oC for 10 min and 60 min are referred to as
‘A1’ and ‘A6’, respectively. The sheets rolled with reduction ratios of 22 % and 41 % right after the annealing at 300 oC for 10 min and 60 min are further designated as ‘A1R2’, ‘A1R4’, ‘A6R2’, and ‘A6R4’, respectively, as listed in Table 1. These clad sheets were sectioned and mechanically polished with a diamond paste with 1 µm size before observing the microstructures of longitudinal-short transverse (L-S) planes by using a scanning electron microscope (SEM, model; JSM-6330F, JEOL, Japan). Phases in the Mg/Al interfacial area were analyzed by using an energy dispersive spectroscopy (EDS). Electron back-scatter diffraction (EBSD) analysis was also conducted by using a field emission scanning electron microscope (FE-SEM, model; Helios NanolabTM, FEI, USA) to obtain inverse pole figure (IPF) and phase maps. Specimens for transmission electron microscopy (TEM) observation were prepared by the focused ion beam (FIB, model: Quanta 3D FEG, FEI, USA) technique. These thin FIB specimens were subsequently used for TEM observations (model; JEM-2100F, JEOL, Japan) under an acceleration voltage of 200 kV. Nano-indentation tests were also conducted on reaction layers formed along the Mg/Al interfaces by using a nano-indenter (model; Nanoindenter XP, MTS, USA) equipped with a triangular Berkovich diamond indenter. All the indentation events were performed at room temperature under a displacement-controlled condition with a constant loading rate of 5 mN/s and the 500 nm depth limit. The load-displacement curves and hardness values of interfacial reaction layers were obtained from this test and triangular indentation marks were observed by a SEM after the test. Plate-type specimens in the longitudinal direction were prepared having the dimensions of 100×10 mm with the thicknesses of 0.4 mm Al ply and 1.5 mm Mg ply for the roller-drum peel tests. They were tested at room temperature under a cross-head speed of 6 mm/min by using a screw-driven universal testing machine (model; 8861, Instron, Canton, MA, USA) of 100 kN capacity. The bonding strength of Mg/Al interface was measured according to the ASTMD3167 specification. After these peel tests, fracture surface of Mg/Al interface was observed and analysed on fifty different locations selected randomly in the same SEM image by using an EDS.
3. Results 3.1. Interfacial Microstructure Backscattered electron images taken from cross-sectional areas of AR (as-roll-bonded), A1 (annealed at 300 oC for 10 min), and A6 (annealed at 300 oC for 60 min) specimens were observed by an SEM, and the results are shown in Figures 1(a) through (c), together with the
EDS composition profiles of Mg and Al in the right side. The Mg/Al interface in AR and A1 specimens is clearly observed from Figures 1(a) and (b) not to contain any pores, cracks, or lateral delamination, as interfacial reaction products are not visible in the Mg/Al interface. As shown in composition profiles of Mg and Al, some atomic diffusion occurs during the roll bonding to form the diffusion layer of 4.7 µm and 5.8 µm in thickness in AR and A1 specimens, respectively. Two kinds of reaction layers are observed in the Mg/Al interface after the annealing at 300 oC for 60 min as shown in Figure 1(c). Backscattered electron images taken from cross-sectional areas of A1R2, A1R4, A6R2, and A6R4 specimens are shown in Figures 2(a) through (d). The Mg/Al interfaces after the one-pass rolling are seen to become curved shape because of considerable amount of deformation at the interface. The reaction layers formed during the prior annealing treatment are maintained in the Mg/Al interface, with reduced thickness after the one-pass rolling. A bright field TEM image and EDS analysis data of reaction layers formed in the Mg/Al interface of AR specimen are shown in Figure 3. Two types of thin reaction layers are observed along the interface as marked by arrows. From the EDS data, microstructures of the points ‘A’ and ‘B’ are identified as γ (Mg17Al12) and β (Mg2Al3) phases, respectively, characteristic intermetallic compounds given in the Mg-Al phase diagram [26]. These reaction layers are thinner than 0.3 µm, too thin to be observed in the SEM micrograph given in Figure 1(a). Bright field TEM images of the interfacial areas of A1 and A6 specimens are shown in Figures 4(a) through (c), together with selected area diffraction (SAD) patterns. The γ (Mg17Al12) and β (Mg2Al3) layers are observed in the Mg and Al alloy sides, respectively, in all of specimens. A1 specimen is shown in Figure 4(a) to exhibit thinly formed in the γ and β layers. Each layer in A6 specimen can be observed to grow up to 2.5 µm in thickness by the annealing for 60 min as shown in Figure 4(b), together with the SAD patterns identifying the layers as γ and β phases. A columnar growth is observed only in the γ layer, but not in the β layer, as shown in Figure 4(c). Bright field TEM images of the interfacial areas of A1R2, A1R4, A6R2, and A6R4 specimens are shown in Figures 5(a) through (d). The γ and β layers in A1R2 and A1R4 specimens are shown in Figures 5(a) and (b) to be thinly formed along the Mg/Al interface, similar to those observed in A1 specimen. The overall thickness of the relatively thick γ and β layers in A6 specimen appears to be reduced after the one-pass rolling as shown in Figures 5(c) and (d) for A6R2 and A6R4 specimens. It is noted here that the columnar growth generated in the γ layer of A6 specimen disappears, and that the overall grain sizes of the γ and β layers are reduced. The average thicknesses of the γ and β layers in the Mg/Al interface were measured from TEM micrographs, and the results are listed in Table 1. These results are also plotted in Figure 6 for a more clear comparison. Very thin layers in AR specimen can be seen to become thicker after the annealing, especially in A6 specimen, in which the layers are grown more rapidly up to 2.0~2.5
µm. The β layer is thicker than the γ layer in all of specimens. The thickness of the γ layer is, however, reduced after the one-pass rolling, but not much change in the β layer thickness in A6 specimen. On the other hand, the γ layer in A1 specimen is grown slightly up to 0.4 µm without changing again the β layer thickness. The overall thickness in A1R2 and A1R4 specimens can also be observed not to change much after the one-pass rolling. 3.2. Nano-indentation Test Results Average hardness values of the γ and β layers as well as the Al and Mg alloys were measured in AR, A6, A6R2, and A6R4 specimens by nano-indentation tests, and the results are shown in Table 1 and Figure 7. The hardness of the γ and β layers is much higher than that of the Al and Mg alloys, and the hardness of the β layer is higher than that of the γ layer. The one-pass rolling raises the hardness of both layers, but the increased amount is much larger in the β layer. Load-displacement curves and SEM micrographs of triangular indentation marks were obtained from indentation tests of the γ and β layers in A6 and A6R4 specimens, and the results are given in Figures 8(a) and (b). The load-displacement curves of the γ and β layers in A6 specimen are not much different with some serrations in both curves as can be observed in Figure 8(a). Peak load of the β layer is, however, much higher than that of the γ layer in A6R4 specimen without any serration in the γ layer. This higher peak load in the β layer appears to match with the higher hardness in the β layer shown in Figure 7. A few micro-cracks are observed in the edge regions of indentation marks in the β layer as indicated by arrows in Figure 8(b). They are also found in the γ layer of A6 specimen, but the number and length are smaller than those in the β layer. Micro-cracks are hardly visible in the γ layer of A6R4 specimen. 3.3. Bonding Strength Bonding strength of the Mg/Al interface was measured from roller-drum peel tests, and the results are shown in Table 1 and Figure 9. AR specimen exhibits an initial bonding strength of 8.7 N/mm. After the annealing treatment, A1 specimen shows twice higher bonding strength compared to that of A6 specimen. This trend is reversed when these annealed specimens are rolled. In other words, the bonding strength of A6 specimen is increased dramatically after the one-pass rolling in comparison to that of A1 specimen exhibiting not much increase. The highest bonding strength is observed in A6R4 specimen. All specimens were observed to peel along the Mg/Al interfacial area after a roller drum peel test. This peeling proceeded by crack initiation and propagation processes along the reaction layers in the Mg/Al interface. SEM micrographs taken from peeled surface of AR, A6, A6R2, and A6R4 specimens are shown in Figures 10(a) through (f). Peeling occurs in a flat mode in AR specimen because of insufficient bonding as observed in Figure 10(a). Peeled surface can be
observed either as the regions of mechanical or metallurgical bonding formed by insufficient or sufficient atomic bonding, respectively, as indicated by arrows in magnified micrograph given in Figure 10(b). The peeled surface in A6 specimen is, on the other hand, observed to show mainly metallurgical bonding region, to exhibit peeling both in the β and γ layers indicated by arrow marks in Figure 10(c), which was confirmed also by subsequent EDS analysis. Here, the peeled surface areas reveal smooth fracture behaviour in the β layer, while cleavage-like fracture mode in the γ layer. Weak river patterns are observed in the cleavage-like area given in Figure 10(d). A6R2 and A6R4 specimens exhibit, on the other hand, fracture facets originated from the β layer without cleavage-like facets related to the γ layer as shown in Figures 10(e) and (f). These smooth facets show somewhat wavy patterns, especially in A6R4 specimen. The cross-sectional area of peeled surface of A6 and A6R4 specimens are shown in Figures 11(a) and (b) to confirm the crack propagation path along the reaction layers in the Mg/Al interface. Considering the crack propagation direction as indicated by arrows in Figure 11(a), the crack passes through both the β and γ layers in A6 specimen, thereby resulting into the peeling both in the β and γ layers. The crack passes through, on the other hand, only the β layer in a rather zig-zag pattern in A6R4 specimen as shown in Figure 11(b). This implies that crack propagation is more difficult in A6R4 specimen compared to that in A6 specimen, because higher energy is continuously required for the crack propagation process.
4. Discussion Bonding properties of metals with an hcp structure including Mg, Cd, and Zr are in general considerably inferior to those of cubic metals such as Al, Cu, and Fe [20, 33, 34]. Most of rollbonded Mg/Al clad sheets are therefore annealed to recover properties of substrate alloys and to improve interfacial bonding properties [29-32]. Reaction layers of γ and β phases are readily formed by diffusion even at relatively low temperatures as expected from the Mg-Al binary phase diagram [25, 26]. They can grow into thicker layers with increasing annealing temperature and time [35-37]. Thermally induced interfacial debonding could, however, occur in thickly grown reaction layers [38]. Most of the recent works on roll-bonded Mg/Al clad sheets have thus been focused on achieving better interfacial bonding strength by restricting the formation of reaction layers or by controlling their thickness [16, 39, 40]. There remain a lot of difficulties in improving interfacial bonding properties simply by changing the annealing time and temperature only. It is therefore essential to control carefully the grain size and morphology of reaction layers as well as their thickness. Diffusion bonding in AR specimen seems to be not sufficient during the roll bonding process. Metallurgical as well as mechanical bonding regions are observed to exist together in AR
specimen as shown in Figure 10(b), resulting into a poor interfacial bonding strength of 8.6 N/mm. These reaction layers in A1 specimen are, on the other hand, seen to grow slightly to expand the region of metallurgical bonding significantly by enhanced interfacial diffusion, resulting into a great improvement of interfacial bonding strength up to 11 N/mm. The regions of mechanical bonding formed partially in the Mg/Al interface could still act as initiation sites for an interfacial embrittlement, which could hardly be removed by an annealing treatment alone. When the annealing time was increased to 60 min, the reaction layers of γ and β grew in thicknesses up to 2.0 µm and 2.7 µm, respectively. The grains in the γ layer are observed to grow in a columnar shape along the vertical direction to the interface as shown in Figures 4(b) and (c). The reason for this columnar growth occurred only in the γ layer can be attributed to the grain size difference between the Mg and Al substrate alloys. The inverse pole figure (IPF) and phase maps of AR specimen obtained from the EBSD analysis are shown in Figures 12(a) and (b), where the interfacial reaction phases are hardly observed in the Mg/Al interface. Though overall grains are coarser in the Al alloy compared to those in the Mg alloy, the grains of the Al alloy can be observed to become much finer near the Mg/Al interface. This appears due to easier recrystallization of the softer Al alloy caused by heavy deformation during the rollbonding process [41, 42]. The difference in grain sizes between the Mg and Al substrate alloys results into the difference in nucleation sites for interfacial reaction phases of γ and β. The nucleation of β phases occurs mostly in the Al alloy side because of refined grain size. In the relatively coarser Mg alloy side, on the other hand, γ phases nucleated at the Mg/Al interface grow continuously along the vertical direction to the interface to form columnar-shape γ grains. This columnar growth of γ grains is previously reported as detrimental to interfacial energy by Jo et al. [43] Peeling readily occurs along the γ layer as well as the β layer during the peel test of A6 specimen, resulting into a large drop in interfacial strength down to 5 N/mm as given in Figure 9. This drop can be recovered by the secondary rolling after the annealing. The thickness of reaction layers in A6 specimen is reduced by the one-pass rolling as observed in A6R2 and A6R4 specimens. The reduced thickness ratios of the γ layer are 26 % and 40 % in A6R2 and A6R4 specimens, respectively, because the hardness of γ phase abruptly decreases above 300 oC [44]. Coarse columnar γ grains are refined into an equi-axed shape by plastic deformation and subsequent recrystallization. Grains in the β layer are also refined, even though the thickness reduction is not that large as given in Figure 6. These changes in thickness together with the size and morphology of grains in the reaction layers are found to greatly influence interfacial properties.
The serrated flow observed for the γ and β layers in A6 specimen implies micro-cracking during nano-indentation tests, which is further confirmed by SEM micrographs of indentation marks given Figure 8(b). Load-displacement curves of the γ and β layers of A6R4 specimen are, however, different from those of A6 specimen. Serrated flow can still be observed in the β layer with reduced trend of micro-cracking behaviour, while the serrations in the γ layer disappear completely without any sign of micro-crack formation. This in turn implies that the γ layer, acted previously as an embrittlement site before the rolling, may not be brittle any longer after the one-pass rolling. Peeled surface given in Figure 10 provides an additional evidence for this, where peeling is observed only along the β layer in A6R2 and A6R4 specimens, but along both the γ and β layers in A6 specimen. The interfacial strength of A6R4 specimen is measured as three times higher than that of A6 specimen. This enhancement could be attributed to the modification of interfacial microstructures and nano-indentation properties after the rolling. Enhancement mechanisms of interfacial bonding strength and microstructural modification by the annealing and secondary rolling treatments are schematically shown in Figure 13. As the annealing time increases, interfacial bonding strength is seen to increase by expanding the regions of metallurgical bonding generated by active atomic diffusion, denoted as a blue solid line in the figure. New brittle phases are, however, formed inside the reaction layers during the annealing to provide detrimental effect on interfacial bonding strength, represented by a green solid line. Reaction layers continue to grow in thickness by sufficient atomic diffusion during the annealing. Thicker reaction layers can easily cause crack initiation and propagation, resulting into the reduction in interfacial bonding strength, now represented by a red solid line in the figure. Furthermore, columnar grain growth in the γ layer critically deteriorates interfacial properties, thereby leading into the lowest bonding strength in A6 specimen. If the one-pass rolling is conducted above a certain reduction ratio after the annealing, interfacial bonding strength could be dramatically improved by microstructural modification as well as thickness reduction of reaction layers, as denoted by a green dotted line. The one-pass rolling can change coarse columnar γ grains into fine equi-axed ones with the thinner γ layer. This microstructural modification of the γ layer appears to play a critical role in toughening and also in reducing the overall thickness of brittle layers. As a result, cracks are initiated and propagated inside the β layer only with wavy patterns, which raises bonding strength dramatically after the rolling, denoted also by a red dotted line in the figure, unlike general expectation of interfacial bonding strength given by a red solid line. SEM micrographs of peeled surface in A6R4 specimen shown in Figures 10(f) and 11(b) seem to support the above explanation of enhanced interfacial bonding strength. Present study on the effect of the one-pass rolling process after the annealing could provide a meaningful method to successfully fabricate Mg/Al clad sheets having excellent interfacial
bonding properties. It can also be useful to understand the interfacial bonding behavior and to suggest optimal annealing and rolling conditions to improve interfacial bonding strength. Twoply Mg/Al clad sheets inevitably contain brittle interfacial reaction layers, which leads to deterioration of interfacial bonding strength by a columnar growth of the γ layer. Interfacial bonding strength is, however, found from present study to improve dramatically after the onepass rolling, which has not been reported so far in any of the previous studies on Mg/Al clad sheets. The improvement of interfacial bonding strength could further be explained by the disappearance of columnar growth and also by grain refinement in the γ layer as observed in A6R4 specimen. This annealing and rolling process appears to provide great advantages in terms of manufacturing, since this one-pass warm-rolling process can utilize the same rolling stand used for the previous roll-bonding process without any additional equipment. Since these Mg/Al clad sheets have outstanding interfacial properties and economic advantages as well, they could provide many new applications to structural materials requiring excellent properties. In order to further enhance microstructures and properties of Mg/Al clad sheets, more intensive studies are currently carried out in relation to new clad alloy designs with enhanced interfacial bonding strength.
5. Conclusions Two-ply Mg/Al clad sheets, fabricated first by a roll-bonding process, were subsequently annealed followed by a one-pass warm rolling. Interfacial bonding properties were evaluated by conducting nano-indentation tests as well as roller-drum peel tests. Interfacial bonding mechanisms were then investigated in relation to the improved interfacial bonding strength. 1) Two types of thin reaction layers formed along the Mg/Al interface were identified as γ (Mg17Al12) and β (Mg2Al3) phases, characteristic intermetallic compounds typically given in the Mg-Al phase diagram. The thickness of reaction layers was increased by extending annealing time. The grains in the γ layer became coarser growing into columnar shape along the vertical direction to the interface during the annealing treatment. This columnar growth occurred only in the γ layer but not in the β layer. 2) After the one-pass rolling, the thickness of the γ layer was reduced but not much change in the β layer thickness. Coarse columnar γ grains became refined into an equi-axed shape by plastic deformation and subsequent recrystallization, while grains in the β layer were refined slightly. These changes in thickness together with the size and morphology of grains in the reaction layers greatly influenced interfacial properties. The serrated flow observed in reaction layers during nano-indentation tests disappeared completely in the γ layer without any sign of
micro-cracks after the one-pass rolling. This in turn implied that the γ layer might not be brittle any longer after the one-pass rolling. 3) The interfacial bonding strength was dramatically improved after the one-pass rolling, which could be attributed due to the disappearance of columnar growth in the γ layer and the grain refinement in the reaction layers. This one-pass rolling process provides great advantages in terms of manufacturing, since it can utilize the same rolling stand used for the previous rollbonding process without any additional equipment. Since these Mg/Al clad sheets have outstanding interfacial properties and economic advantages as well, they can provide various new applications to structural materials requiring excellent properties.
Acknowledgments This study was supported by a grant (10037273) from the Fundamental R&D Program for Core Technology of Materials funded by the Ministry of Knowledge Economy, Korea, and by POSCO under the contract No. 2014Y008.
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Table 1. Thickness and hardness of interfacial layers and bonding strength of Mg/Al interface in the two-ply clad sheets.
Annealing Reduction Time Ratio Specimen at of 573 K Rolling (300 oC)
Thickness of γ Layer (µm)
Thickness Thickness of of Total β Layer Interfacial (µm) Layer (µm)
Hardness of γ Layer (GPa)
Hardness of β Layer (GPa)
Bonding Strength of Mg/Al Interface (N/mm)
AR
-
-
0.16±0.02
0.12±0.01
0.28±0.03
-
-
8.67±0.59
A1
10 min
-
0.39±0.03
0.16±0.03
0.55±0.13
-
-
10.22±0.60
A6
60 min
-
1.99±0.09
2.69±0.29
4.69±0.34
1.54±0.02
1.65±0.12
5.10±1.10
A1R2
10 min
22 %
0.35±0.02
0.17±0.01
0.51±0.02
-
-
10.83±1.86
A1R4
10 min
41 %
0.40±0.06
0.14±0.02
0.54±0.05
-
-
12.02±1.33
A6R2
60 min
22 %
1.47±0.10
2.59±0.14
4.06±0.08
2.12±0.13
2.65±0.22
15.47±1.43
A6R4
60 min
41 %
1.19±0.17
2.57±0.16
3.76±0.29
2.17±0.19
2.74±0.01
18.01±0.68
Fig. 1 Back-scattered electron images taken from cross-sectional areas of (a) AR, (b) A1, and (c) A6 specimens. Mg/Al interfaces are clearly visible without any pores, cracks, or lateral delamination. Fig. 2 Backscattered electron images taken from cross-sectional areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. Reaction layers are found in the Mg/Al interface. Fig. 3 Bright field TEM image and EDS analysis data of reaction layers formed in the Mg/Al interface of AR specimen. Microstructures of points ‘A’ and ‘B’ are identified as γ (Mg17Al12) and β (Mg2Al3) phases, respectively, from the EDS data. Fig. 4 Bright field TEM images and selected area diffraction (SAD) patterns of the Mg/Al interfacial areas of (a) A1, and (b) and (c) A6 specimens. The γ (Mg17Al12) and β (Mg2Al3) layers are observed in the Mg and Al alloy sides, respectively. Fig. 5 Bright field TEM images of the Mg/Al interfacial areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. The γ and β layers are thinly formed along the Mg/Al interface in A1R2 and A1R4 specimens. The overall thickness of the relatively thick γ and β layers in A6 specimen is reduced after the one-pass rolling as shown in (c) and (d) for A6R2 and A6R4 specimens. Fig. 6 Average thickness of the γ and β layers in the Mg/Al interface of AR, A1, A6, A1R2, A1R4, A6R2, and A6R4 specimens. Fig. 7 Average hardness of the γ and β layers as well as the Al and Mg alloys in AR, A6, A6R2, and A6R4 specimens. Fig. 8 (a) Load-displacement curves and (b) SEM micrographs of triangular indentation marks obtained from indentation tests of the γ and β layers in A6 and A6R4 specimens. A few microcracks are observed in the edge regions of indentation marks. Fig. 9 Bonding strength of the Mg/Al interface measured from roller-drum peel test of AR, A6, A6R2, and A6R4 specimens. The bonding strength of A6 specimen is increased dramatically after the one-pass rolling in comparison to that of A1 specimen exhibiting not much increase. Fig. 10 SEM fractographs taken from peeled surface of (a) and (b) AR, (c) and (d) A6, (e) A6R2, and (f) A6R4 specimens. Fig. 11 SEM micrographs taken from cross-sectional area of peeled surface in (a) A6 and (b) A6R4 specimens. Crack passes through the β and γ layers in A6 specimen, while through the β layer only in a rather zig-zag pattern in A6R4 specimen.
Fig. 12 (a) Inverse pole figure (IPF) map and (b) phase map of AR specimen. Grains in the Al alloy are refined near the Mg/Al interface due to recrystallization of the softer Al alloy caused by heavy plastic deformation during the roll-bonding process. Fig. 13 Schematic diagram illustrating enhancement mechanisms of interfacial bonding strength and microstructural modification by the annealing and secondary one-pass rolling.
Figure(s)
Fig. 1 Back-scattered electron images taken from cross-sectional areas of (a) AR, (b) A1, and (c) A6 specimens. Mg/Al interfaces are clearly visible without any pores, cracks, or lateral delamination.
Fig. 2 Backscattered electron images taken from cross-sectional areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. Reaction layers are found in the Mg/Al interface.
Fig. 3 Bright field TEM image and EDS analysis data of reaction layers formed in the Mg/Al interface of AR specimen. Microstructures of points ‘A’ and ‘B’ are identified as γ (Mg17Al12) and β (Mg2Al3) phases, respectively, from the EDS data.
AR
Mg
A
Mg17Al12 (A)
A
[At%]
Mg
59.83
Al
40.17
Al B
Mg2Al3 (B)
1ঙ
B
[At%]
Mg
41.10
Al
61.40
Fig. 4 Bright field TEM images and selected area diffraction (SAD) patterns of the Mg/Al interfacial areas of (a) A1, and (b) and (c) A6 specimens. The γ (Mg17Al12) and β (Mg2Al3) layers are observed in the Mg and Al alloy sides, respectively.
Fig. 5 Bright field TEM images of the Mg/Al interfacial areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. The γ and β layers are thinly formed along the Mg/Al interface in A1R2 and A1R4 specimens. The overall thickness of the relatively thick γ and β layers in A6 specimen is reduced after the one-pass rolling as shown in (c) and (d) for A6R2 and A6R4 specimens.
Fig. 6 Average thickness of the γ and β layers in the Mg/Al interface of AR, A1, A6, A1R2, A1R4, A6R2, and A6R4 specimens.
Average Thickness (mm)
5 4 A6
3
g b (g+b)
2
A1
1
g b (g+b)
0 AR
A6 A1
A6R2 A1R2
Specimen
A6R4 A1R4
Fig. 7 Average hardness of the γ and β layers as well as the Al and Mg alloys in AR, A6, A6R2, and A6R4 specimens.
Average Hardness (GPa)
3.0 2.5 2.0
β
1.5
γ
1.0
Mg 0.5
Al
0.0 AR
A6
Specimen
A6R2
A6R4
Fig. 8 (a) Load-displacement curves and (b) SEM micrographs of triangular indentation marks obtained from indentation tests of the γ and β layers in A6 and A6R4 specimens. A few micro-cracks are observed in the edge regions of indentation marks.
Fig. 9 Bonding strength of the Mg/Al interface measured from roller-drum peel test of AR, A6, A6R2, and A6R4 specimens. The bonding strength of A6 specimen is increased dramatically after the one-pass rolling in comparison to that of A1 specimen exhibiting not much increase.
20
A6R4
Bonding Strength (N/mm)
18
A6R2
16 14 12 10
A1 AR
A1R4
8
A1R2
6 4
A6
2 0 AR
A6 A1
A6R2 A1R2
Specimen
A6R4 A1R4
Fig. 10 SEM fractographs taken from peeled surface of (a) and (b) AR, (c) and (d) A6, (e) A6R2, and (f) A6R4 specimens.
Fig. 11 SEM micrographs taken from cross-sectional area of peeled surface in (a) A6 and (b) A6R4 specimens. Crack passes through the β and γ layers in A6 specimen, while through the β layer only in a rather zig-zag pattern in A6R4 specimen.
Fig. 12 (a) Inverse pole figure (IPF) map and (b) phase map of AR specimen. Grains in the Al alloy are refined near the Mg/Al interface due to recrystallization of the softer Al alloy caused by heavy plastic deformation during the roll-bonding process.
Fig. 13 Schematic diagram illustrating enhancement mechanisms of interfacial bonding strength and microstructural modification by the annealing and secondary one-pass rolling.
Improvement of interfacial bonding strength in roll-bonded Mg/Al clad Sheets through annealing and secondary rolling process Jung-Su Kim, Kwang Seok Lee, Yong Nam Kwon, Byeong-Joo Lee, Young Won Chang, Sunghak Lee
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PII: DOI: Reference:
S0921-5093(15)00051-9 http://dx.doi.org/10.1016/j.msea.2015.01.035 MSA31973
To appear in:
Materials Science & Engineering A
Received date: 30 October 2014 Revised date: 13 January 2015 Accepted date: 14 January 2015 Cite this article as: Jung-Su Kim, Kwang Seok Lee, Yong Nam Kwon, ByeongJoo Lee, Young Won Chang, Sunghak Lee, Improvement of interfacial bonding strength in roll-bonded Mg/Al clad Sheets through annealing and secondary rolling process, Materials Science & Engineering A, http://dx.doi.org/10.1016/j. msea.2015.01.035 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Improvement of Interfacial Bonding Strength in Roll-Bonded Mg/Al Clad Sheets through Annealing and Secondary Rolling Process Jung-Su Kima, Kwang Seok Leeb, Yong Nam Kwonb, Byeong-Joo Leea, Young Won Changc, and Sunghak Leea,* a
Center for Advanced Aerospace Materials
Pohang University of Science and Technology, Pohang, 790-784, Korea b
Materials Deformation Department, Light Metal Division
Korea Institute of Materials Science, Changwon, 641-010, Korea c
Graduate Institute of Ferrous Technology
Pohang University of Science and Technology, Pohang 790-784, Korea +82-54-279-2140(Telephone) / +82-54-279-5887(Fax) *Corresponding Author e-mail: [email protected] (S. Lee)
Abstract The effect of annealing and secondary rolling process has been investigated in relation to evolution of interface microstructure as well as interfacial bonding mechanisms of roll-bonded two-ply Mg/Al clad sheets in this study. Two types of thin reaction layers, viz., γ (Mg17Al12) and β (Mg2Al3) phase layers, were observed to form along the Mg/Al interface, the thickness of which was found to increase with annealing time. The grains in the γ layer were grown coarsely in a columnar shape, as they grew continuously in coarse-grained Mg substrate alloy side. The thickness of the γ layer was found to reduce significantly after the secondary one-pass rolling. At the same time, the initially columnar γ grains became refined into equi-axed grains by recrystallization caused by large plastic deformation during the secondary rolling. The serrated flow observed in the reaction layers during a nano-indentation test in the γ layer disappeared completely without any sign of micro-cracks after the secondary rolling. This in turn implies that the γ layer, acted as an embrittlement site before the rolling, may not be brittle any longer after the rolling. The grain refinement and disappearance of columnar grains in the γ layer appear to greatly improve the interfacial bonding strength in these clad sheets. This secondary rolling process seems to provide a great advantage in terms of manufacturing, since the same rolling stand for roll-bonding process can also be used without any additional equipment. Keywords: Mg/Al clad sheet, annealing, secondary one-pass rolling, interfacial bonding, γ (Mg17Al12) phase, β (Mg2Al3) phase
1. Introduction Multi-layered metallic clad sheets combining two or more alloys have recently been developed to provide a new multi-functional property utilizing characteristic properties of each metallic alloy. Disadvantages of one alloy sheet can therefore be effectively complemented or improved to provide entirely new application areas in various industries such as automotive, aerospace, and kitchen utensils industries, and so on [1-5]. These multi-layered clad sheets have generally been fabricated by using Al alloys, Ti alloys, or stainless steels to utilize their excellent heat conductivity and corrosion resistance as well as other specific mechanical properties [6, 7]. Global warming has especially become a serious issue all over the world including carbon dioxide emissions to generate a significant attention in structural materials field. In order to reduce carbon dioxide emission, a number of research and development efforts have actively been conducted on lightweight materials such as Mg and Al alloys [8-10]. Mg alloys are the lightest-weight metal commercially available today, having various advantageous properties in terms of specific strength, damping capacity, electromagnetic shielding, and dimensional stability [11, 12]. Actual industrial application areas of Mg alloys are, however, quite limited due to their inherent drawbacks in relation to poor corrosion resistance, formability, and surface quality. In order to expand the applications of Mg alloys, various types of clad Mg alloys have been developed by bonding with other alloys such as stainless steels and Al alloys to complement the drawbacks of Mg alloys [13-16]. Direct solid-state bonding between Mg and Al alloys is, however, seriously restricted due to limited formability of Mg alloys at low and intermediate temperatures [15]. Various solid-state bonding processes of Mg/Al sheets have been developed to date including roll-bonding, friction-bonding, diffusion-bonding, and explosion-bonding [17-23]. Among these, roll-bonding is the most widely used process for fabricating Mg/Al clad sheets, considering economic advantages in terms of convenient utilization of conventional rolling facilities, productivity, and easy combination of clad metals [24]. A warm roll-bonding is generally used to circumvent the poor room-temperature formability of Mg alloys followed by subsequently annealing treatment to control mechanical properties of clad sheets. Brittle reaction phases are, however, inevitably formed along the solid-state-bonded interface in the form of intermetallic compounds, viz., γ (Mg17Al12) and β (Mg2Al3) phases, as readily expected from the Mg-Al binary phase diagram [25, 26], to deteriorate overall mechanical properties as well as interfacial bonding strength. These brittle reaction layers have recently been reported by Lee et al. [27] to act as interfacial delamination sites during the mid-stage flexural bending test of a three-ply laminated composite of stainless steel/Al/Mg. A thin interlayer of pure Zn was inserted between Mg and Al alloy sheets to prevent this interfacial delamination problem as suggested by Zhao et
al. [28] The insertion of a high-purity Zn foil was found to retard or suppress the formation of γ and β phases during the solid-state diffusion bonding. This method may not be suitable for the roll-bonding process of wide-width clad sheets in view of manufacturing cost. The size, morphology, and thickness of reaction phases should be optimally controlled in order to improve the interfacial bonding strength of Mg/Al clad sheets. The effects of these parameters on interfacial bonding strength have not, however, been clarified to date, except the generally accepted fact that interfacial bonding strength could be enhanced by controlling the thickness of reaction phases to a certain level [29-32]. Elucidation of mechanisms in relation to microstructural modification is required as well, since the only variation of thickness of reaction layer cannot provide a plausible explanation on interfacial bonding strength [16]. In addition, a cost effective fabrication method of Mg/Al clad sheets should be developed by utilizing conventional facilities such as rolling stands to obtain adequate microstructures and required properties simultaneously. This fabrication idea is quite useful especially in relation to the production costs of Mg/Al clad sheets having excellent microstructures and properties, taking full advantages of individual Mg and Al alloys. Mg/Al clad sheets were therefore fabricated in this study by a roll-bonding process together with a subsequent annealing treatment followed by a warm rolling. This process route provided the merits of economical fabrication of Mg/Al clad sheets and also the improvement of their interfacial bonding strength simultaneously. Interfacial microstructures of fabricated sheets were then analyzed, and their properties were evaluated by conducting nano-indentation as well as roller-drum peel tests. Interfacial bonding mechanisms were also investigated in relation to the improving interfacial bonding strength, by focusing specifically on how the size, morphology, and thickness of reaction phases affected the initiation and propagation of cracks in the Mg/Al interface.
2. Experimental procedures Multi-pass warm-rolling was performed first using sheets of 3004 Al and AZ31 Mg alloys at 400 oC with the reduction ratio of 60 % to fabricate 3.3-mm-thick two-ply clad sheets. The 3004 Al alloy sheet had the thickness of 1.0 mm with the composition of Al-(1.0~1.5)Mn(0.8~1.3)Mg-(≤0.7)Fe-(≤0.3)Si, while the AZ31 Mg alloy sheet had the thickness of 2.3 mm with the composition of Mg-(2.5~3.5)Mn-(0.6~1.4)Zn-(0.2~1.0Mn)-(≤0.1)Si. These roll-bonded clad sheets were then subsequently annealed at 300 oC for 10 min and 60 min before air-cooling. Right after this annealing treatment, one-pass warm-rolling (300 oC) was carried out for some of the sheets with the total reduction ratio of 22 % or 41 %. The as-roll-bonded clad sheet is designated as ‘AR’, while the sheets annealed at 300 oC for 10 min and 60 min are referred to as
‘A1’ and ‘A6’, respectively. The sheets rolled with reduction ratios of 22 % and 41 % right after the annealing at 300 oC for 10 min and 60 min are further designated as ‘A1R2’, ‘A1R4’, ‘A6R2’, and ‘A6R4’, respectively, as listed in Table 1. These clad sheets were sectioned and mechanically polished with a diamond paste with 1 µm size before observing the microstructures of longitudinal-short transverse (L-S) planes by using a scanning electron microscope (SEM, model; JSM-6330F, JEOL, Japan). Phases in the Mg/Al interfacial area were analyzed by using an energy dispersive spectroscopy (EDS). Electron back-scatter diffraction (EBSD) analysis was also conducted by using a field emission scanning electron microscope (FE-SEM, model; Helios NanolabTM, FEI, USA) to obtain inverse pole figure (IPF) and phase maps. Specimens for transmission electron microscopy (TEM) observation were prepared by the focused ion beam (FIB, model: Quanta 3D FEG, FEI, USA) technique. These thin FIB specimens were subsequently used for TEM observations (model; JEM-2100F, JEOL, Japan) under an acceleration voltage of 200 kV. Nano-indentation tests were also conducted on reaction layers formed along the Mg/Al interfaces by using a nano-indenter (model; Nanoindenter XP, MTS, USA) equipped with a triangular Berkovich diamond indenter. All the indentation events were performed at room temperature under a displacement-controlled condition with a constant loading rate of 5 mN/s and the 500 nm depth limit. The load-displacement curves and hardness values of interfacial reaction layers were obtained from this test and triangular indentation marks were observed by a SEM after the test. Plate-type specimens in the longitudinal direction were prepared having the dimensions of 100×10 mm with the thicknesses of 0.4 mm Al ply and 1.5 mm Mg ply for the roller-drum peel tests. They were tested at room temperature under a cross-head speed of 6 mm/min by using a screw-driven universal testing machine (model; 8861, Instron, Canton, MA, USA) of 100 kN capacity. The bonding strength of Mg/Al interface was measured according to the ASTMD3167 specification. After these peel tests, fracture surface of Mg/Al interface was observed and analysed on fifty different locations selected randomly in the same SEM image by using an EDS.
3. Results 3.1. Interfacial Microstructure Backscattered electron images taken from cross-sectional areas of AR (as-roll-bonded), A1 (annealed at 300 oC for 10 min), and A6 (annealed at 300 oC for 60 min) specimens were observed by an SEM, and the results are shown in Figures 1(a) through (c), together with the
EDS composition profiles of Mg and Al in the right side. The Mg/Al interface in AR and A1 specimens is clearly observed from Figures 1(a) and (b) not to contain any pores, cracks, or lateral delamination, as interfacial reaction products are not visible in the Mg/Al interface. As shown in composition profiles of Mg and Al, some atomic diffusion occurs during the roll bonding to form the diffusion layer of 4.7 µm and 5.8 µm in thickness in AR and A1 specimens, respectively. Two kinds of reaction layers are observed in the Mg/Al interface after the annealing at 300 oC for 60 min as shown in Figure 1(c). Backscattered electron images taken from cross-sectional areas of A1R2, A1R4, A6R2, and A6R4 specimens are shown in Figures 2(a) through (d). The Mg/Al interfaces after the one-pass rolling are seen to become curved shape because of considerable amount of deformation at the interface. The reaction layers formed during the prior annealing treatment are maintained in the Mg/Al interface, with reduced thickness after the one-pass rolling. A bright field TEM image and EDS analysis data of reaction layers formed in the Mg/Al interface of AR specimen are shown in Figure 3. Two types of thin reaction layers are observed along the interface as marked by arrows. From the EDS data, microstructures of the points ‘A’ and ‘B’ are identified as γ (Mg17Al12) and β (Mg2Al3) phases, respectively, characteristic intermetallic compounds given in the Mg-Al phase diagram [26]. These reaction layers are thinner than 0.3 µm, too thin to be observed in the SEM micrograph given in Figure 1(a). Bright field TEM images of the interfacial areas of A1 and A6 specimens are shown in Figures 4(a) through (c), together with selected area diffraction (SAD) patterns. The γ (Mg17Al12) and β (Mg2Al3) layers are observed in the Mg and Al alloy sides, respectively, in all of specimens. A1 specimen is shown in Figure 4(a) to exhibit thinly formed in the γ and β layers. Each layer in A6 specimen can be observed to grow up to 2.5 µm in thickness by the annealing for 60 min as shown in Figure 4(b), together with the SAD patterns identifying the layers as γ and β phases. A columnar growth is observed only in the γ layer, but not in the β layer, as shown in Figure 4(c). Bright field TEM images of the interfacial areas of A1R2, A1R4, A6R2, and A6R4 specimens are shown in Figures 5(a) through (d). The γ and β layers in A1R2 and A1R4 specimens are shown in Figures 5(a) and (b) to be thinly formed along the Mg/Al interface, similar to those observed in A1 specimen. The overall thickness of the relatively thick γ and β layers in A6 specimen appears to be reduced after the one-pass rolling as shown in Figures 5(c) and (d) for A6R2 and A6R4 specimens. It is noted here that the columnar growth generated in the γ layer of A6 specimen disappears, and that the overall grain sizes of the γ and β layers are reduced. The average thicknesses of the γ and β layers in the Mg/Al interface were measured from TEM micrographs, and the results are listed in Table 1. These results are also plotted in Figure 6 for a more clear comparison. Very thin layers in AR specimen can be seen to become thicker after the annealing, especially in A6 specimen, in which the layers are grown more rapidly up to 2.0~2.5
µm. The β layer is thicker than the γ layer in all of specimens. The thickness of the γ layer is, however, reduced after the one-pass rolling, but not much change in the β layer thickness in A6 specimen. On the other hand, the γ layer in A1 specimen is grown slightly up to 0.4 µm without changing again the β layer thickness. The overall thickness in A1R2 and A1R4 specimens can also be observed not to change much after the one-pass rolling. 3.2. Nano-indentation Test Results Average hardness values of the γ and β layers as well as the Al and Mg alloys were measured in AR, A6, A6R2, and A6R4 specimens by nano-indentation tests, and the results are shown in Table 1 and Figure 7. The hardness of the γ and β layers is much higher than that of the Al and Mg alloys, and the hardness of the β layer is higher than that of the γ layer. The one-pass rolling raises the hardness of both layers, but the increased amount is much larger in the β layer. Load-displacement curves and SEM micrographs of triangular indentation marks were obtained from indentation tests of the γ and β layers in A6 and A6R4 specimens, and the results are given in Figures 8(a) and (b). The load-displacement curves of the γ and β layers in A6 specimen are not much different with some serrations in both curves as can be observed in Figure 8(a). Peak load of the β layer is, however, much higher than that of the γ layer in A6R4 specimen without any serration in the γ layer. This higher peak load in the β layer appears to match with the higher hardness in the β layer shown in Figure 7. A few micro-cracks are observed in the edge regions of indentation marks in the β layer as indicated by arrows in Figure 8(b). They are also found in the γ layer of A6 specimen, but the number and length are smaller than those in the β layer. Micro-cracks are hardly visible in the γ layer of A6R4 specimen. 3.3. Bonding Strength Bonding strength of the Mg/Al interface was measured from roller-drum peel tests, and the results are shown in Table 1 and Figure 9. AR specimen exhibits an initial bonding strength of 8.7 N/mm. After the annealing treatment, A1 specimen shows twice higher bonding strength compared to that of A6 specimen. This trend is reversed when these annealed specimens are rolled. In other words, the bonding strength of A6 specimen is increased dramatically after the one-pass rolling in comparison to that of A1 specimen exhibiting not much increase. The highest bonding strength is observed in A6R4 specimen. All specimens were observed to peel along the Mg/Al interfacial area after a roller drum peel test. This peeling proceeded by crack initiation and propagation processes along the reaction layers in the Mg/Al interface. SEM micrographs taken from peeled surface of AR, A6, A6R2, and A6R4 specimens are shown in Figures 10(a) through (f). Peeling occurs in a flat mode in AR specimen because of insufficient bonding as observed in Figure 10(a). Peeled surface can be
observed either as the regions of mechanical or metallurgical bonding formed by insufficient or sufficient atomic bonding, respectively, as indicated by arrows in magnified micrograph given in Figure 10(b). The peeled surface in A6 specimen is, on the other hand, observed to show mainly metallurgical bonding region, to exhibit peeling both in the β and γ layers indicated by arrow marks in Figure 10(c), which was confirmed also by subsequent EDS analysis. Here, the peeled surface areas reveal smooth fracture behaviour in the β layer, while cleavage-like fracture mode in the γ layer. Weak river patterns are observed in the cleavage-like area given in Figure 10(d). A6R2 and A6R4 specimens exhibit, on the other hand, fracture facets originated from the β layer without cleavage-like facets related to the γ layer as shown in Figures 10(e) and (f). These smooth facets show somewhat wavy patterns, especially in A6R4 specimen. The cross-sectional area of peeled surface of A6 and A6R4 specimens are shown in Figures 11(a) and (b) to confirm the crack propagation path along the reaction layers in the Mg/Al interface. Considering the crack propagation direction as indicated by arrows in Figure 11(a), the crack passes through both the β and γ layers in A6 specimen, thereby resulting into the peeling both in the β and γ layers. The crack passes through, on the other hand, only the β layer in a rather zig-zag pattern in A6R4 specimen as shown in Figure 11(b). This implies that crack propagation is more difficult in A6R4 specimen compared to that in A6 specimen, because higher energy is continuously required for the crack propagation process.
4. Discussion Bonding properties of metals with an hcp structure including Mg, Cd, and Zr are in general considerably inferior to those of cubic metals such as Al, Cu, and Fe [20, 33, 34]. Most of rollbonded Mg/Al clad sheets are therefore annealed to recover properties of substrate alloys and to improve interfacial bonding properties [29-32]. Reaction layers of γ and β phases are readily formed by diffusion even at relatively low temperatures as expected from the Mg-Al binary phase diagram [25, 26]. They can grow into thicker layers with increasing annealing temperature and time [35-37]. Thermally induced interfacial debonding could, however, occur in thickly grown reaction layers [38]. Most of the recent works on roll-bonded Mg/Al clad sheets have thus been focused on achieving better interfacial bonding strength by restricting the formation of reaction layers or by controlling their thickness [16, 39, 40]. There remain a lot of difficulties in improving interfacial bonding properties simply by changing the annealing time and temperature only. It is therefore essential to control carefully the grain size and morphology of reaction layers as well as their thickness. Diffusion bonding in AR specimen seems to be not sufficient during the roll bonding process. Metallurgical as well as mechanical bonding regions are observed to exist together in AR
specimen as shown in Figure 10(b), resulting into a poor interfacial bonding strength of 8.6 N/mm. These reaction layers in A1 specimen are, on the other hand, seen to grow slightly to expand the region of metallurgical bonding significantly by enhanced interfacial diffusion, resulting into a great improvement of interfacial bonding strength up to 11 N/mm. The regions of mechanical bonding formed partially in the Mg/Al interface could still act as initiation sites for an interfacial embrittlement, which could hardly be removed by an annealing treatment alone. When the annealing time was increased to 60 min, the reaction layers of γ and β grew in thicknesses up to 2.0 µm and 2.7 µm, respectively. The grains in the γ layer are observed to grow in a columnar shape along the vertical direction to the interface as shown in Figures 4(b) and (c). The reason for this columnar growth occurred only in the γ layer can be attributed to the grain size difference between the Mg and Al substrate alloys. The inverse pole figure (IPF) and phase maps of AR specimen obtained from the EBSD analysis are shown in Figures 12(a) and (b), where the interfacial reaction phases are hardly observed in the Mg/Al interface. Though overall grains are coarser in the Al alloy compared to those in the Mg alloy, the grains of the Al alloy can be observed to become much finer near the Mg/Al interface. This appears due to easier recrystallization of the softer Al alloy caused by heavy deformation during the rollbonding process [41, 42]. The difference in grain sizes between the Mg and Al substrate alloys results into the difference in nucleation sites for interfacial reaction phases of γ and β. The nucleation of β phases occurs mostly in the Al alloy side because of refined grain size. In the relatively coarser Mg alloy side, on the other hand, γ phases nucleated at the Mg/Al interface grow continuously along the vertical direction to the interface to form columnar-shape γ grains. This columnar growth of γ grains is previously reported as detrimental to interfacial energy by Jo et al. [43] Peeling readily occurs along the γ layer as well as the β layer during the peel test of A6 specimen, resulting into a large drop in interfacial strength down to 5 N/mm as given in Figure 9. This drop can be recovered by the secondary rolling after the annealing. The thickness of reaction layers in A6 specimen is reduced by the one-pass rolling as observed in A6R2 and A6R4 specimens. The reduced thickness ratios of the γ layer are 26 % and 40 % in A6R2 and A6R4 specimens, respectively, because the hardness of γ phase abruptly decreases above 300 oC [44]. Coarse columnar γ grains are refined into an equi-axed shape by plastic deformation and subsequent recrystallization. Grains in the β layer are also refined, even though the thickness reduction is not that large as given in Figure 6. These changes in thickness together with the size and morphology of grains in the reaction layers are found to greatly influence interfacial properties.
The serrated flow observed for the γ and β layers in A6 specimen implies micro-cracking during nano-indentation tests, which is further confirmed by SEM micrographs of indentation marks given Figure 8(b). Load-displacement curves of the γ and β layers of A6R4 specimen are, however, different from those of A6 specimen. Serrated flow can still be observed in the β layer with reduced trend of micro-cracking behaviour, while the serrations in the γ layer disappear completely without any sign of micro-crack formation. This in turn implies that the γ layer, acted previously as an embrittlement site before the rolling, may not be brittle any longer after the one-pass rolling. Peeled surface given in Figure 10 provides an additional evidence for this, where peeling is observed only along the β layer in A6R2 and A6R4 specimens, but along both the γ and β layers in A6 specimen. The interfacial strength of A6R4 specimen is measured as three times higher than that of A6 specimen. This enhancement could be attributed to the modification of interfacial microstructures and nano-indentation properties after the rolling. Enhancement mechanisms of interfacial bonding strength and microstructural modification by the annealing and secondary rolling treatments are schematically shown in Figure 13. As the annealing time increases, interfacial bonding strength is seen to increase by expanding the regions of metallurgical bonding generated by active atomic diffusion, denoted as a blue solid line in the figure. New brittle phases are, however, formed inside the reaction layers during the annealing to provide detrimental effect on interfacial bonding strength, represented by a green solid line. Reaction layers continue to grow in thickness by sufficient atomic diffusion during the annealing. Thicker reaction layers can easily cause crack initiation and propagation, resulting into the reduction in interfacial bonding strength, now represented by a red solid line in the figure. Furthermore, columnar grain growth in the γ layer critically deteriorates interfacial properties, thereby leading into the lowest bonding strength in A6 specimen. If the one-pass rolling is conducted above a certain reduction ratio after the annealing, interfacial bonding strength could be dramatically improved by microstructural modification as well as thickness reduction of reaction layers, as denoted by a green dotted line. The one-pass rolling can change coarse columnar γ grains into fine equi-axed ones with the thinner γ layer. This microstructural modification of the γ layer appears to play a critical role in toughening and also in reducing the overall thickness of brittle layers. As a result, cracks are initiated and propagated inside the β layer only with wavy patterns, which raises bonding strength dramatically after the rolling, denoted also by a red dotted line in the figure, unlike general expectation of interfacial bonding strength given by a red solid line. SEM micrographs of peeled surface in A6R4 specimen shown in Figures 10(f) and 11(b) seem to support the above explanation of enhanced interfacial bonding strength. Present study on the effect of the one-pass rolling process after the annealing could provide a meaningful method to successfully fabricate Mg/Al clad sheets having excellent interfacial
bonding properties. It can also be useful to understand the interfacial bonding behavior and to suggest optimal annealing and rolling conditions to improve interfacial bonding strength. Twoply Mg/Al clad sheets inevitably contain brittle interfacial reaction layers, which leads to deterioration of interfacial bonding strength by a columnar growth of the γ layer. Interfacial bonding strength is, however, found from present study to improve dramatically after the onepass rolling, which has not been reported so far in any of the previous studies on Mg/Al clad sheets. The improvement of interfacial bonding strength could further be explained by the disappearance of columnar growth and also by grain refinement in the γ layer as observed in A6R4 specimen. This annealing and rolling process appears to provide great advantages in terms of manufacturing, since this one-pass warm-rolling process can utilize the same rolling stand used for the previous roll-bonding process without any additional equipment. Since these Mg/Al clad sheets have outstanding interfacial properties and economic advantages as well, they could provide many new applications to structural materials requiring excellent properties. In order to further enhance microstructures and properties of Mg/Al clad sheets, more intensive studies are currently carried out in relation to new clad alloy designs with enhanced interfacial bonding strength.
5. Conclusions Two-ply Mg/Al clad sheets, fabricated first by a roll-bonding process, were subsequently annealed followed by a one-pass warm rolling. Interfacial bonding properties were evaluated by conducting nano-indentation tests as well as roller-drum peel tests. Interfacial bonding mechanisms were then investigated in relation to the improved interfacial bonding strength. 1) Two types of thin reaction layers formed along the Mg/Al interface were identified as γ (Mg17Al12) and β (Mg2Al3) phases, characteristic intermetallic compounds typically given in the Mg-Al phase diagram. The thickness of reaction layers was increased by extending annealing time. The grains in the γ layer became coarser growing into columnar shape along the vertical direction to the interface during the annealing treatment. This columnar growth occurred only in the γ layer but not in the β layer. 2) After the one-pass rolling, the thickness of the γ layer was reduced but not much change in the β layer thickness. Coarse columnar γ grains became refined into an equi-axed shape by plastic deformation and subsequent recrystallization, while grains in the β layer were refined slightly. These changes in thickness together with the size and morphology of grains in the reaction layers greatly influenced interfacial properties. The serrated flow observed in reaction layers during nano-indentation tests disappeared completely in the γ layer without any sign of
micro-cracks after the one-pass rolling. This in turn implied that the γ layer might not be brittle any longer after the one-pass rolling. 3) The interfacial bonding strength was dramatically improved after the one-pass rolling, which could be attributed due to the disappearance of columnar growth in the γ layer and the grain refinement in the reaction layers. This one-pass rolling process provides great advantages in terms of manufacturing, since it can utilize the same rolling stand used for the previous rollbonding process without any additional equipment. Since these Mg/Al clad sheets have outstanding interfacial properties and economic advantages as well, they can provide various new applications to structural materials requiring excellent properties.
Acknowledgments This study was supported by a grant (10037273) from the Fundamental R&D Program for Core Technology of Materials funded by the Ministry of Knowledge Economy, Korea, and by POSCO under the contract No. 2014Y008.
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Table 1. Thickness and hardness of interfacial layers and bonding strength of Mg/Al interface in the two-ply clad sheets.
Annealing Reduction Time Ratio Specimen at of 573 K Rolling (300 oC)
Thickness of γ Layer (µm)
Thickness Thickness of of Total β Layer Interfacial (µm) Layer (µm)
Hardness of γ Layer (GPa)
Hardness of β Layer (GPa)
Bonding Strength of Mg/Al Interface (N/mm)
AR
-
-
0.16±0.02
0.12±0.01
0.28±0.03
-
-
8.67±0.59
A1
10 min
-
0.39±0.03
0.16±0.03
0.55±0.13
-
-
10.22±0.60
A6
60 min
-
1.99±0.09
2.69±0.29
4.69±0.34
1.54±0.02
1.65±0.12
5.10±1.10
A1R2
10 min
22 %
0.35±0.02
0.17±0.01
0.51±0.02
-
-
10.83±1.86
A1R4
10 min
41 %
0.40±0.06
0.14±0.02
0.54±0.05
-
-
12.02±1.33
A6R2
60 min
22 %
1.47±0.10
2.59±0.14
4.06±0.08
2.12±0.13
2.65±0.22
15.47±1.43
A6R4
60 min
41 %
1.19±0.17
2.57±0.16
3.76±0.29
2.17±0.19
2.74±0.01
18.01±0.68
Fig. 1 Back-scattered electron images taken from cross-sectional areas of (a) AR, (b) A1, and (c) A6 specimens. Mg/Al interfaces are clearly visible without any pores, cracks, or lateral delamination. Fig. 2 Backscattered electron images taken from cross-sectional areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. Reaction layers are found in the Mg/Al interface. Fig. 3 Bright field TEM image and EDS analysis data of reaction layers formed in the Mg/Al interface of AR specimen. Microstructures of points ‘A’ and ‘B’ are identified as γ (Mg17Al12) and β (Mg2Al3) phases, respectively, from the EDS data. Fig. 4 Bright field TEM images and selected area diffraction (SAD) patterns of the Mg/Al interfacial areas of (a) A1, and (b) and (c) A6 specimens. The γ (Mg17Al12) and β (Mg2Al3) layers are observed in the Mg and Al alloy sides, respectively. Fig. 5 Bright field TEM images of the Mg/Al interfacial areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. The γ and β layers are thinly formed along the Mg/Al interface in A1R2 and A1R4 specimens. The overall thickness of the relatively thick γ and β layers in A6 specimen is reduced after the one-pass rolling as shown in (c) and (d) for A6R2 and A6R4 specimens. Fig. 6 Average thickness of the γ and β layers in the Mg/Al interface of AR, A1, A6, A1R2, A1R4, A6R2, and A6R4 specimens. Fig. 7 Average hardness of the γ and β layers as well as the Al and Mg alloys in AR, A6, A6R2, and A6R4 specimens. Fig. 8 (a) Load-displacement curves and (b) SEM micrographs of triangular indentation marks obtained from indentation tests of the γ and β layers in A6 and A6R4 specimens. A few microcracks are observed in the edge regions of indentation marks. Fig. 9 Bonding strength of the Mg/Al interface measured from roller-drum peel test of AR, A6, A6R2, and A6R4 specimens. The bonding strength of A6 specimen is increased dramatically after the one-pass rolling in comparison to that of A1 specimen exhibiting not much increase. Fig. 10 SEM fractographs taken from peeled surface of (a) and (b) AR, (c) and (d) A6, (e) A6R2, and (f) A6R4 specimens. Fig. 11 SEM micrographs taken from cross-sectional area of peeled surface in (a) A6 and (b) A6R4 specimens. Crack passes through the β and γ layers in A6 specimen, while through the β layer only in a rather zig-zag pattern in A6R4 specimen.
Fig. 12 (a) Inverse pole figure (IPF) map and (b) phase map of AR specimen. Grains in the Al alloy are refined near the Mg/Al interface due to recrystallization of the softer Al alloy caused by heavy plastic deformation during the roll-bonding process. Fig. 13 Schematic diagram illustrating enhancement mechanisms of interfacial bonding strength and microstructural modification by the annealing and secondary one-pass rolling.
Figure(s)
Fig. 1 Back-scattered electron images taken from cross-sectional areas of (a) AR, (b) A1, and (c) A6 specimens. Mg/Al interfaces are clearly visible without any pores, cracks, or lateral delamination.
Fig. 2 Backscattered electron images taken from cross-sectional areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. Reaction layers are found in the Mg/Al interface.
Fig. 3 Bright field TEM image and EDS analysis data of reaction layers formed in the Mg/Al interface of AR specimen. Microstructures of points ‘A’ and ‘B’ are identified as γ (Mg17Al12) and β (Mg2Al3) phases, respectively, from the EDS data.
AR
Mg
A
Mg17Al12 (A)
A
[At%]
Mg
59.83
Al
40.17
Al B
Mg2Al3 (B)
1ঙ
B
[At%]
Mg
41.10
Al
61.40
Fig. 4 Bright field TEM images and selected area diffraction (SAD) patterns of the Mg/Al interfacial areas of (a) A1, and (b) and (c) A6 specimens. The γ (Mg17Al12) and β (Mg2Al3) layers are observed in the Mg and Al alloy sides, respectively.
Fig. 5 Bright field TEM images of the Mg/Al interfacial areas of (a) A1R2, (b) A1R4, (c) A6R2, and (d) A6R4 specimens. The γ and β layers are thinly formed along the Mg/Al interface in A1R2 and A1R4 specimens. The overall thickness of the relatively thick γ and β layers in A6 specimen is reduced after the one-pass rolling as shown in (c) and (d) for A6R2 and A6R4 specimens.
Fig. 6 Average thickness of the γ and β layers in the Mg/Al interface of AR, A1, A6, A1R2, A1R4, A6R2, and A6R4 specimens.
Average Thickness (mm)
5 4 A6
3
g b (g+b)
2
A1
1
g b (g+b)
0 AR
A6 A1
A6R2 A1R2
Specimen
A6R4 A1R4
Fig. 7 Average hardness of the γ and β layers as well as the Al and Mg alloys in AR, A6, A6R2, and A6R4 specimens.
Average Hardness (GPa)
3.0 2.5 2.0
β
1.5
γ
1.0
Mg 0.5
Al
0.0 AR
A6
Specimen
A6R2
A6R4
Fig. 8 (a) Load-displacement curves and (b) SEM micrographs of triangular indentation marks obtained from indentation tests of the γ and β layers in A6 and A6R4 specimens. A few micro-cracks are observed in the edge regions of indentation marks.
Fig. 9 Bonding strength of the Mg/Al interface measured from roller-drum peel test of AR, A6, A6R2, and A6R4 specimens. The bonding strength of A6 specimen is increased dramatically after the one-pass rolling in comparison to that of A1 specimen exhibiting not much increase.
20
A6R4
Bonding Strength (N/mm)
18
A6R2
16 14 12 10
A1 AR
A1R4
8
A1R2
6 4
A6
2 0 AR
A6 A1
A6R2 A1R2
Specimen
A6R4 A1R4
Fig. 10 SEM fractographs taken from peeled surface of (a) and (b) AR, (c) and (d) A6, (e) A6R2, and (f) A6R4 specimens.
Fig. 11 SEM micrographs taken from cross-sectional area of peeled surface in (a) A6 and (b) A6R4 specimens. Crack passes through the β and γ layers in A6 specimen, while through the β layer only in a rather zig-zag pattern in A6R4 specimen.
Fig. 12 (a) Inverse pole figure (IPF) map and (b) phase map of AR specimen. Grains in the Al alloy are refined near the Mg/Al interface due to recrystallization of the softer Al alloy caused by heavy plastic deformation during the roll-bonding process.
Fig. 13 Schematic diagram illustrating enhancement mechanisms of interfacial bonding strength and microstructural modification by the annealing and secondary one-pass rolling.