Fabrication and mechanical properties of ultrafine structured dissimilar laminated metal composite sheets (LMCS)

1 Introduction

In recent years, more and more interest has been focused on ultrafine grained (UFG) multilayered sheets consisting of alternant bimetal systems fabricated by accumulative roll bonding (ARB), such as Al/Al(Sr) [1], Al/Cu [2], Al/Ni [3, 4], Al/Zn [5], Ti/Al [6], Cu/Zr [7], Cu/Ni [8], and Al/Al alloys [9–11]. ARB is an emerging severe plastic deformation (SPD) technology; it can obviously improve the mechanical properties of materials. This persistently strong push to research on these multilayered sheets is caused by twofold effects. One is that the ultrafine grain or subgrain structure combined with the high density of lattice defects for SPD metals leads to the high yield strength [12–16]. The other is that the intrinsic characteristic of multilayered architecture has its own advantages, including the comparably light weight, improved impact and fracture toughness, excellent wear resistance and damping behavior, and other compromised properties [17, 18]. However, only a limited range of metals is suitable for obtaining bulk UFG structure by ARB. The remarkable grain refinement associated with densely cumulated dislocations could usually be realized only below the recrystallization temperature of the materials [19]. Some processed metals with poor formability may not achieve such a severe cold plastic deformation or provide a sound bonding interface. Furthermore, the requirements of a large load and powerful rolling mills by the ARB process are also obstacles to their industrial applications. Therefore, some endeavors by using just plain rolling thickness reduction regime to refine the grains down to ultrafine structure have been proposed, but most of them were carried out on monolithic metals, for example, duplex ferrite-martensite steel in which the grain refinement was accelerated by the strain concentration due to complex plastic flow of hard and soft phases [20] and pure metals or alloys with considerable deformability so that the ultrafine structure could be obtained by cryogenic rolling [20–25] or high-ratio differential speed rolling [26–29]. In these cases, the formation of UFG was considered to attribute to the suppression of dynamic recovery and the intense shear strain induced during rolling.

The objective of this study is to develop a modified route to produce bulk ultrafine structured Al alloy laminated composite sheets without severe cold plastic deformation. The fabrication of laminated metal composite sheet (LMCS) included the following four processes: (1) hot roll bonding of five-layer composite sheet, (2) ARB at elevated temperature to fabricate the multilayered sheet, (3) recrystallization annealing and the subsequent cold rolling at room temperature (RT) to obtain an ultrafine structure and, (4) low-temperature annealing to optimize the mechanical properties. A 1280 layer ultrafine structure Al alloy composite sheet was synthesized successfully by this process. The mechanical properties and microstructure of this LMCS are investigated in this paper. The accelerated grain refinement and related mechanisms are also analyzed according to the microstructure observation.

2 Materials and methods

Two extreme Al series, commercial pure 1100 and high-strength 7075 Al alloys, were selected and used in this study; the chemical compositions are shown in Table 1. The two Al alloys, which are widely applied in modern industry, have great differences in mechanical properties and formability.

The fabrication experiment and microstructure characterization of 1100/7075 LMCS were carried out in Material Processing Laboratory and Center of Electron Microscopy of Chongqing University. After being degreased and wire-brushed, three pieces of Al 1100 and two pieces of Al 7075 alloys were stacked alternately to form a five-layered composite slab of 12-mm thickness, held at 460°C for 30 min in a preheated furnace, and then hot roll-bonded with 50% thickness reduction by one pass. The roll bonding process of five-layered composite sheet and its rolling schedule are shown as section I in Figure 1. The obtained 6-mm-thick five-layered composite sheet was then cut in half for degreasing, wire-brushing, and stacking again to form a 12-mm-thick 10-layered composite sheet and hot roll-bonded in the same process as above, which is referred to as the first ARB cycle (denoted as ARB-1). After ARB for eight cycles and an immediate thickness reduction from 6 to 4 mm by one pass, a superior roll-bonded 1280 layer composite sheet 4 mm thick was fabricated. The hot ARB process and its rolling schedule of LMCS are shown as section II in Figure 1. This as-rolled 1280 layer composite sheet was then annealed at 400°C for 1 h to improve its deformability. This coarse-grained sample was then cold rolled (CR) to 1 mm thick at RT by several passes with about 10% thickness reduction per pass (i.e., total CR thickness reduction 75%, equivalent strain 1.6). The cold rolling process of 1280 layer composite sheet and its rolling schedule are shown as section III in Figure 1. For comparison, a five-layered composite sheet was also obtained in the same way. The cold rolling was executed under lubricated conditions on a rolling mill with a roll diameter of 170 mm and rolling speed of 0.2 m/s. In order to optimize the mechanical properties, some CR samples were annealed (150°C for 30 min, 200°C for 5 min) or solid solution (SS) peak aging treated (SS at 475°C for 1 h, immediately quenched in water, and then aged at 120°C for 72 h).

Figure 1

Schematic diagram of cold rolling process and its rolling schedule of LMCS.

The microstructure of the processed LMCS was investigated using optical micrographs, scanning electron microscopy (SEM, FEI Nova 400, Chongqing University, Chongqing, China) equipped with energy-dispersive X-ray spectroscopy (EDS, Oxford Instruments, Chongqing University, Chongqing, China), and transmission electron microscopy (TEM, Zeiss Libra 200 FE, Chongqing University, Chongqing, China) in the transverse direction (TD) of the samples. Tensile tests were carried out at RT on the AG-X testing machine operated at a constant crosshead speed with an initial strain rate of 5.5×10^-4/s. The gauge section of the tensile specimen was 15 mm long and 7 mm wide, and the tensile direction was parallel to the rolling direction (RD) of the sample.

3 Results

The thickness of the metals decreased with the increase in ARB cycles. However, the dislocation density of the final composite sheet cannot increase during the rolling process. The reason is that the laminated metal composite sheets were held for 30 min at 460°C before the ARB process was performed, and the process annealing treatment was executed at 400°C for 1 h before cold rolling process was done. Both processes will result in the recrystallization and grain boundary migration and neutralize the variation of dislocation density and the strengthening action of subdivision of grain.

Figure 2 shows the micrographs of 1280 layer composite sheets before and after cold rolling. Figure 2A shows the optical micrograph of 4-mm-thick 1280 layer composite sheet that was annealed at 400°C for 1 h before cold rolling. From Figure 2A, the laminated structure and coarse grains can be observed clearly; the average sizes of the grain are ∼42 μm length and ∼12 μm width. Theoretically, the average thickness of each layer metal is about 5 μm. Therefore, a single grain can necessarily cross several adjacent 1100 and 7075 layers. Good metallurgical bonding was realized between two Al series metals by the ARB in conjunction with the intermediate heat treatment. Figure 2B shows the scanning electron microscope- backscattered electron (SEM-BSE) micrograph of the processed 1280 layer composite sheet with 1-mm thickness. Cold rolling with 75% reduction resulted in a wavy lamellar microstructure including some discontinuous oxide layers along the RD (Figure 2B). The corrugated microstructure with local layer directions deviating from the RD was caused mainly by the incompatible plastic flow of the adjacent 1100 and 7075 layers during the rolling process.

Figure 2

Micrographs of TD section of the processed 1280 layer composite sheets: (A) optical image for 4-mm thick 1280 layer sheet after annealing at 400°C for 1 h; (B) SEM micrograph of 1-mm thick 1280 layer sheet after 75% cold rolling, showing wavy lamellar microstructure with some oxide layers and constituent particles, which were identified as Al₇Cu₂Fe by EDS analysis.

Quadir et al. [30] suggested that the oxide layer could also be fractured into discontinuous ones by the operation of shear bands, which could be further validated in Figure 3. Figure 3A shows the larger magnification of shear bands and grain distortions around the oxides. Because of the ultrafine component layers with an average thickness of ∼1 μm, both the oxide layers and large constituent particles would remarkably influence the local flows and distributions of layers, which revealed a complex strain distribution throughout the whole composite sheet during cold rolling process. Figure 3B shows the schematic diagram of grain refinement and shear directions around the oxides. There are strong shear actions between the alloy layers near the oxides due to the poor formability of oxides. Except for the rolling reduction, additional shear actions promote the grain refinement of metals near the oxides.

Figure 3

Shear bands and grain distortions around the oxides. (A) SEM-BSE larger magnification showing shear bands and grain distortions around the oxides; (B) schematic diagram of grain refinement and shear bands around the oxides.

The microstructure of annealed 4-mm thick 1280 layer composite sheet is typical mix grain structure, as Figure 2A shows. There are small grains in the 7075 layer around the large grain. It is easy to result in complexity stress and strain distribution during rolling reduction, and occurring stress concentrates in the grain zone with different scales, as Figure 4A shows. The local stress concentration will obviously increase the local stress, promote the increase of dislocation density at this zone, and accelerate the refinement of grains. The large grains were refined and divided into many small grains under shear actions and compression actions. The accelerating grain refinement mechanism is illustrated in Figure 4B.

Figure 4

Schematic diagram showing grain refinement during cold rolling. (A) Grain distribution of multilayer composite sheet; (B) grain refinement mechanism.

Figure 5A shows elongated ultrafine grains and subgrains (average size ∼300 nm) along the layer direction in both alternant 1100 and 7075 layers after 75% CR. There is some limited structure growth (average size ∼450 nm) after annealing at 150°C for 30 min. This shows that recovery had occurred (Figure 5B). The zones of two different metals are distinguished and marked based on the distributions of dispersoids and precipitates in the 7075 alloy. It can be observed that there are obviously necking and shear characteristics.

Figure 5

SEM-BSE micrographs of TD section of the processed 1280 layer composite sheets (A) showing ultrafine structure in the layers after 75% cold rolling, (B) showing coarsened structure after 75% cold rolling and then annealing at 150°C for 30 min. The particles (50–100 nm in size) distributed in 7075 layers were identified as the equilibrium precipitate MgZn₂ by EDS analysis [31].

More profound investigation of the 7075 layer is shown in Figure 6. The 7075 layer of the CR 1280 layer composite sheet revealed more refined structure including some network structure within well-defined grains (120–400 nm in size) and subgrains with thick cell walls (200–300 nm in size), which corresponded to a higher density of dislocations (Figure 6A). The selected area electron diffraction (SAED) pattern in Figure 6A was ring-like, suggesting the existence of large misorientations. Moreover, the rod-like Al₁₂Mg₂Cr phase is typically dispersoid, having a size of about 20–500 nm in the 7075 Al alloy. Once the phase formed, it could not be completely dissolved at high temperature [32], but could not be found in Figure 2A. As the Al₁₂Mg₂Cr dispersoids cannot be dissolved to an appreciable extent by solid-state thermal treatments [33], their “disappearance” from the 7075 layer of the CR 1280 layer composite sheet was considered to be attributed to their significant fragmentation into smaller particles, much higher density of dislocations, and associated intense strain field. Annealing at 150°C for 30 min after 75% CR led to thinner cell walls and diminished dislocations in the 7075 layer of 1280 layer composite sheet (Figure 6B). The SAED pattern in Figure 6B was still ring-like but with less clustered diffraction spots, which suggested a lower fraction of low-angle grain boundaries.

Figure 6

TEM micrographs of the TD section in the 7075 layer of the processed samples: (A) 1280 layer composite sheet of 75% cold rolling; (B) 1280 layer composite sheet after 75% cold rolling and then annealing at 150°C for 30 min. Both the SAED patterns in (A) and (B) were taken from a ∼1.2-μm-diameter sample area.

Figure 7 shows the representative tensile engineering stress-strain curves of the processed samples. The tensile properties obtained from these curves are summarized in Table 2 in which a comparison with that of related monolithic metals previously published in references is also listed. The CR 1280 layer composite sheet exhibited higher strength by 30% and slightly larger elongation than the five-layered composite sheet with the same CR thickness reduction of 75% (curves a and e). As the ratio of the volume fraction between 1100 and 7075 Al alloy was kept unchanged at 1:1 during the whole rolling process, the higher strength was considered to be derived from the more refined structure, higher density of dislocations, and more dispersoids in the 1280 layer composite sheet (Figure 6A), as will be discussed later. It is noteworthy in Table 2 that the tensile strength of this CR 1280 layer 1100/7075 composite sheet could be even comparable to that of 7075 Al alloy sheet subjected to ARB for five cycles at 250°C [34], which highlighted the crucial role of rolling temperature on strengthening metals by ARB.

Figure 7

Tensile engineering stress-strain curves of the processed composite sheets.

An optimization of strength and ductility for the 1280 layer composite sheet was achieved by annealing at 150°C for 30 min, with almost twice the elongation and only 9% yield strength loss compared with that of the CR temper (curves a and b). In order to further evaluate this optimized strength, a 7075 Al alloy sheet was cold rolled from the original 3 mm to 1 mm thick, and it was found that they had similar tensile strength (curves b and f). Annealing at 200°C for 5 min gave a reasonably acceptable ductility (8.9% total elongation and 6.6% uniform elongation) but at the sacrifice of too much strength (16% yield strength loss, curve c). The conventional SS-peak aging treatment revealed a broad strain hardening with about 24% uniform elongation, but the yield strength was low (curve d), which was less than half of that of the CR temper. Many previous research works on SPD metals pointed out a popular strength relationship of two to four times between the UFG structure and its coarse-grained counterpart [12, 13, 34], with which the yield strength comparison between the CR temper and the SS temper of this 1280 layer composite sheet may agree reasonably well if the contribution of the precipitates strengthening in 7075 layers were to be eliminated. It is also interesting to note that the whole tensile strain hardening of the annealed 1280 layer composite sheet was not consecutive or smooth (curves b and c); instead it experienced several plateaus (i.e., almost zero strain hardening rate) by steps with an abrupt stress increase between the two adjacent plateaus, as will also be discussed later.

4 Discussion

Comparing with the five-layered composite sheet, although much more strain (equivalent strain 6.4) is introduced into the 1280 layer composite sheet during the ARB cycles at 460°C, which can refine the layers, they may be little effective in making a difference on dislocation density between these two samples because of the dynamic recrystallization during the ARB process, the grain boundary migrations during the repeated preheating periods at 460°C, and the subsequent annealing at 400°C for 1 h. The more refined structure and higher density of dislocations for the 1280 layer composite sheets are mainly derived from the cold rolling process. First, the additional shear strain in the vicinity of the layer interfaces is caused by the drag and friction between the adjacent 1100 and 7075 layers due to their incompatible plastic flow during the cold rolling process, which can facilitate the accumulation of dislocations. As the layers became ultrafine, this more frequent distribution of redundant shear strain throughout the whole sheet associated with the rolling strain contributed to a higher total strain, which can accelerate the grain subdivision and refinement [3, 4, 9–11]. Second, more oxide layers are introduced into the 1280 layer composite sheet during hot roll bonding cycles and can be further fractured into fragments or particles by the cold rolling conducted afterward. As for the small oxide particles, besides interaction with dislocations, they also act as dislocation sources to accelerate the dislocation nucleation by dislocations bowing around these particles and leaving dislocation loops behind, both of which would lead to the rapid accumulation of dislocations [2, 9–11, 36, 37]. Recent research on the effect of coarse constituent particles on deformation behaviors of 3104 Al alloy validated significant local distortions and increase of misorientation angles near the particles [38], which may resemble the situation of this study, that is, the large oxide fragments that are sporadically embedded in the layers (Figure 3A and B) also played an important role in grain refinement.

In previous research works on the correlation between strength and microstructure for deformed metals [14–16, 39], there are generally two strengthening mechanisms proposed: grain boundary strengthening caused by relatively high angle boundaries with a classical Hall-Petch relationship, and forest dislocation strengthening caused by the dislocations in very low angle boundaries and in the volume between the boundaries. However, in our study for multilayered sheet with different number of layers, besides the above two strengthening mechanisms, a relationship of yield strength-layer thickness similar to Hall-Petch law is also established, indicating that the layer interface may also be one of the dominant strengthening factors, because these interfaces act as barriers to dislocation glide, especially in the case of ultrafine layers. Moreover, as the small alumina particles are extruded from the interfaces into the layers and the Al₁₂Mg₂Cr dispersoids are also fragmented into smaller particles during the rolling process, another dispersion strengthening (i.e., Orowan mechanism) therefore occurred. However, this strengthening contribution is supposed to be minimal because of the negligibly effective dispersoids.

Generally, the enhanced ductility of the annealed 1280 layer composite sheets should be attributed to the recovery accompanied by the dislocation annihilation and rearrangement and the internal stress relief, but here two other mechanisms are proposed to elucidate the characteristic strain hardening of the annealed 1280 layer composite sheets during tensile straining (Figure 7, curves b and c). The first is related to the observed shear bands during tensile deformation, as shown in Figure 8. Figure 8A shows the SEM-BSE micrograph of the gauge section of the 150°C annealed 1280 layer composite sheet after tensile to fracture. Figure 8B shows the shear direction, fluctuant bonding interface, and local necking of LMCS. The propagation of a shear band involved a process of repeated stress concentration and relaxation. It initially formed and propagated in the soft 1100 layer due to the local strain concentration, which is a stress relaxation process; as it propagated to the layer interface, the propagation is inhibited by the hard 7075 layer, which caused a local stress concentration, and also more stresses are needed to stimulate the shear band across the interface. Statistically, a balance between these stress concentrations and relaxations from many shear bands contributed to the almost zero strain hardening rate on the individual plateau, while the break of this balance by more shear bands propagating across the layer interfaces and penetrating into the 7075 layers accompanied the abrupt stress increase. This is because the 7075 layers are more resistant against the shear bands propagating across, and therefore more stresses are needed for further straining. Second, as there are more than 1000 layer interfaces through just 1-mm thickness, the average space of these interfaces has been calculated statistically to be ∼1 μm, i.e., very high average interface area per unit volume S_V with an estimated value of 10⁶/m; therefore, the dynamic recovery is accelerated by these dense interfaces where dislocations slip and sink easily, which lowered the rate of dislocation accumulation during tensile straining. It should also be pointed out that both of the mechanisms are based on the already existing recovered state after annealing as the prerequisite, as in order to activate them, the dislocation storage and slip during tensile straining are indispensable, which need extra room in between the dislocation networks. This is why the stepwise strain hardening process is not pronounced for the CR 1280 layer composite sheet during tensile test (Figure 7, curve a).

Figure 8

SEM-BSE micrograph on gauge section of the 150°C annealed 1280 layer composite sheet after tensile to fracture (corresponding to curve b in Figure 7) showing several shear bands: (A) SEM-BSE magnification showing shear bands; (B) schematic diagram of shear directions at each layer.

5 Conclusions

In summary, we have successfully fabricated the bulk ultrafine structured 1100/7075 Al alloy LMCS by rolling. It is believed that this strategy can also be extended to other aluminum alloys series, bimetal or multimetal systems with poor deformability below the recrystallization temperature for synthesizing bulk ultrafine structured multilayered composite sheets. The additional shear actions, which are caused by the incompatibility of deformation dissimilar materials, promote the refinement of grains. The stress concentration, which is caused by complex crystals of metals, accelerates the partition of large grains and fragments of dispersoids during cold rolling process. These actions work together and result in the ultrafine structured dissimilar LMCS. In the present study, although the fine layers, ultrafine structure, increased dislocation density, and abundant dispersoids in the CR 1280 layer composite sheet contributed to its higher strength by 30% than that of the five-layered composite sheet; the ductility is improved at the expense of 9% yield strength loss by annealing at 150°C for 30 min. Although the yield strength decreased, the SS-peak aging treatment revealed a broad strain hardening with about 24% uniform elongation. If we execute a thermomechanical treatment, which can take full advantage of the precipitate behaviors of the 7075 Al alloy, to substitute for the intermediate recrystallization annealing and the final low-temperature annealing, the strength and ductility of dissimilar LMCS will be possibly enhanced simultaneously.