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Article

Effects of Cold Rolling on the Microstructure and Properties of Al/TiB2 Laminated Composites Fabricated by Accumulative Roll Bonding

by
Wenchao Sun
,
Zhilei Xiang
,
Jihao Li
,
Zian Yang
,
Yang Han
and
Ziyong Chen
*
College of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, China
*
Author to whom correspondence should be addressed.
Materials 2026, 19(5), 1031; https://doi.org/10.3390/ma19051031
Submission received: 12 February 2026 / Revised: 2 March 2026 / Accepted: 6 March 2026 / Published: 8 March 2026
(This article belongs to the Section Metals and Alloys)
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Abstract

Al/TiB2 aluminum alloy laminates were fabricated using a combination of accumulative roll bonding (ARB) and cold rolling processes. The Al/TiB2 interface and microstructure were meticulously characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). The mechanical properties of the laminates were assessed through tensile testing. The experimental results demonstrate that with an increasing cold rolling reduction, a laminated composite sheet with a nanocrystalline structure was successfully produced. The critical strain for the onset of plastic instability was also investigated. The findings indicate that as the cold rolling reduction increases, severe necking occurs in the Al12Zn2.2Mg1.7Cu3TiB2 layer. At a reduction of 80%, the necking region approaches fracture. Tensile results reveal that this pronounced necking has a detrimental effect on the strength of the laminate. It is proposed that the plastic instability originates from shear bands, and the mechanical property mismatch between the constituent layers is identified as the primary reason for the localized preferential deformation.

1. Introduction

In recent years, severe plastic deformation (SPD) techniques have been widely employed in the fabrication of aluminum matrix composites [1,2,3]. These processes are capable of refining the microstructure of alloys and composites to ultrafine-grained levels, leading to significant improvements in material properties. Commonly used SPD methods currently include equal-channel angular pressing (ECAP), high-pressure torsion (HPT), multi-axial forging, and accumulative roll bonding (ARB) [4,5,6]. Among these various SPD techniques, the ARB process demonstrates distinct advantages. Its primary benefits lie in its compatibility with conventional rolling equipment and its capability to produce large-scale sheet materials, making it particularly suitable for industrial applications [2].
The Accumulative Roll Bonding (ARB) process was first proposed by Y. Saito [4] in 1998. As a significant severe plastic deformation (SPD) technique, it is commonly employed to achieve grain refinement and enhance the bonding quality of sheets through multi-step solid-state bonding [2]. In a typical procedure, two or more sheets of identical dimensions are stacked, fastened with rivets, and then roll-bonded. The process involves rolling the layered stack with a predetermined reduction, followed by cutting the bonded sheet into two halves, which are then stacked again and subjected to a subsequent rolling cycle. A key advantage of ARB is its ability to accumulate a high strain while maintaining the original dimensions of the workpiece. Compared to other SPD techniques such as equal-channel angular pressing and high-pressure torsion, ARB is regarded as a highly promising method due to its simplicity in tooling requirements, high production efficiency, and unique capability for continuous production of large-scale sheets. To date, ARB has been successfully applied to fabricate and study a wide range of bimetallic systems (e.g., Al/Cu [7], Al/Ni [8], Ti/Al [9], Zn/Sn [10], Cu/Zn [11]) and even some trimetallic systems (e.g., Al/Cu/Sn [12], Al/Ti/Mg [13], Ti/Al/Nb [14]). Research in this field primarily focuses on analyzing the evolution of the interfacial structure and on manipulating the microstructure of the constituent layers, with the aim of understanding their consequent effects on the mechanical properties of the composite material. Ling Ou et al. [15] fabricated 6061Al/Cu composites using the accumulative roll bonding (ARB) process, achieving a maximum tensile strength of 613 MPa. The results indicate that as the strain rate increases, the necking parameters of the Al and Cu layers gradually decrease, while the interfacial shear strength increases. Ultimately, at a strain rate of 30 s−1, the tensile strength reaches its peak value of 613 MPa, with hardness values of 186 HV for the Al layer and 203 HV for the Cu layer, respectively.
Utilizing the characteristics of the ARB process, such as high strain rates, significant shear strain, changes in strain path, and the low thermal conductivity of commercially pure titanium, Yang et al. [9] successfully fabricated Ti/Al composite sheets with grain sizes of approximately 200–300 nm. With increasing strain, the formation of an ultrafine-grained structure and the fragmentation of the Ti layers contributed to a continuous enhancement of the sheet’s strength and hardness. M. Askarpour et al. [16] fabricated an ultrafine-grained Al-TiB2 laminated composite via the accumulative roll bonding (ARB) process. After seven ARB cycles, the in situ TiB2 particle-reinforced aluminum composite developed an ultrafine-grained structure with an average grain size of approximately 300 nm. In a study on Al/Cu composite laminates, Eizadjou et al. [7] investigated the microstructural evolution and mechanical properties after different ARB cycles. They observed that as the number of ARB cycles increased, the Cu layers progressively fragmented and became uniformly distributed as reinforcing phases within the Al matrix. This led to significant improvements in both strength and ductility. The strength enhancement was primarily attributed to strain hardening in the initial deformation stages and grain refinement strengthening in the later stages. However, a common challenge in processing laminated composites via ARB is the occurrence of plastic instability, often resulting from the differing flow and plastic properties of the constituent layers. For instance, in their research on AA1050/AA6061 composite sheets, Su et al. [17] reported that an additional shear effect, induced by the property mismatch between the constituents, led to plastic instability within the harder layer.
During an investigation into the microstructural evolution and mechanical properties of Zn/Sn composite sheets, Mashhadi et al. [10] observed that pronounced necking occurred in the Zn layer after the second ARB cycle, which led to a significant deterioration in the mechanical performance of the laminate. The primary reason is that the onset of necking compromises the ability of the harder layer to bear the applied load, thereby adversely affecting the overall properties. In research aimed at optimizing against plastic instability, Zeng et al. [18] successfully fabricated large-scale Cu/Ta laminated composites by employing an ARB process combined with intermediate annealing [12,13]. They demonstrated that processing at elevated temperatures can effectively suppress the occurrence of plastic instability. However, it was also found that the mechanical properties of the laminated composites processed under these high-temperature conditions were generally moderate [19]. Consequently, a key focus of current research is to develop strategies that effectively inhibit plastic instability while simultaneously maintaining high mechanical performance in the composite sheets.
In this study, laminated composites were fabricated from an aluminum alloy and an aluminum matrix composite using the accumulative roll bonding (ARB) process. Subsequent cold rolling was employed to successfully produce sheets featuring a nano/ultrafine-grained structure and excellent mechanical properties. The evolution of the microstructure with increasing cold rolling reduction was systematically investigated. Furthermore, the mechanical properties of the composites under various reductions were comprehensively evaluated based on detailed microstructural analysis.

2. Materials and Methods

This study employed two types of self-developed aluminum alloys with nominal compositions of Al12Zn2.2Mg1.7Cu and Al12Zn2.2Mg1.7Cu3TiB2. The as-received materials, with dimensions of 200 mm × 60 mm × 4 mm, were used as the starting blanks. Both sheets were fabricated via hot extrusion at 410 °C. Additionally, the micron/submicron TiB2 particles used in this study were synthesized as Al/TiB2 grain refiners via an oxide-based reaction method, employing aluminum powder, titanium powder, titanium oxide, and boric acid as the raw reactants. The actual measured chemical compositions of the two base materials are provided in Table 1.
The ARB process was conducted using the two aforementioned aluminum alloy sheets, each with an initial thickness of 4 mm. Rolling was performed on a two-high rolling mill with a roll diameter of 600 mm, a roll face width of 500 mm, and a maximum rolling force of 60 tons. The rolling speed was set at 0.2 m/s, and no lubricant was applied between the rolls and the workpiece surfaces. Figure 1 is a picture of the rolling mill equipment (Yaxin General Machinery Co, Ltd., Wuxi, China) used in the experiment. Prior to rolling, the faying surfaces of the sheets were subjected to a treatment sequence involving degreasing, scratch brushing with a steel wire brush, and cleaning with alcohol. This procedure was essential to completely remove oxide films and work-hardened layers while generating a specific surface roughness, thereby promoting mutual atomic diffusion and the extrusion of base metal across the interface to ensure sound bonding. Subsequently, two sheets of Alloy A and one sheet of Alloy B were stacked to form an A/B/A “sandwich” configuration. To prevent interlayer sliding during rolling, the stack was securely fastened along one side using rivets.
Step 1: The initial three-layer “sandwich” stack with a total thickness of 12 mm was homogenized at 460 °C for 20 min, followed by hot rolling in a single pass to a thickness of 6 mm. This initial state is designated as 0 ARB cycles. Step 2: The resulting composite plate from Step 1 was sectioned longitudinally into two halves. After surface treatment (cleaning and scratch brushing), the two sections were stacked anew, riveted, and reheated to 460 °C for 20 min. The reassembled pack was then hot-rolled with a 50% reduction to a final thickness of 6 mm, completing the first ARB cycle. Step 2 was repeated iteratively until a total of 6 ARB cycles were accomplished. This procedure resulted in laminated composites with a progressively increasing number of layers: 3, 6, 12, 24, 48, and 96 layers, corresponding to 0 to 6 cycles, respectively. The TiB2 content in the obtained laminated composite sheet is 1 wt.%. The entire experimental process was conducted in an ambient laboratory environment without a protective atmosphere. A schematic diagram illustrating the ARB process for fabricating the multilayer composite sheets is presented in Figure 2.
The primary objective of this study is to achieve an ultrafine-grained microstructure and enhance the mechanical properties of Al/TiB2 laminated composites fabricated by accumulative roll bonding (ARB) through room-temperature cold rolling deformation. Additionally, the critical strain for the onset of severe plastic instability is determined. The process involved five passes with thickness reductions of 16%, 32%, 48%, 64%, and 80%, respectively. The microstructure of the composites was characterized using a Quanta FEG 650 scanning electron microscope (SEM, FEI Company, Hillsboro, OR, USA) and a JEOL 2100 transmission electron microscope (TEM, JEOL Ltd, Tokyo, Japan). The SEM system was equipped with an Oxford X-MAXN-80 energy-dispersive X-ray spectroscopy (EDS, FEI Company, Hillsboro, OR, USA) detector and an electron backscatter diffraction (EBSD, FEI Company, Hillsboro, OR, USA) detector. For EBSD analysis, the samples were prepared by mechanical polishing followed by electrolytic polishing. TEM samples were first ground mechanically to a thickness below 70 µm, punched into 3-mm discs, and then thinned using a dual-jet electro-polisher with a solution of 30% nitric acid and 70% methanol at 248 K until perforation. The JEOL 2100 TEM (JEOL Ltd., Tokyo, Japan) was equipped with super-EDS, a high-angle annular dark-field (HAADF) detector, a bright-field detector, and two dark-field detectors. High-resolution TEM (HRTEM) images were acquired in TEM mode using a Thermo Scientific Velox system (Thermo Fisher Scientific, Newark, DE, USA) and analyzed with Gatan Microscopy Suite (GMS-3, Gatan, Inc., Pleasanton, CA, USA) software. Tensile tests were conducted on an MTS 810 testing machine (MTS Systems, Eden Prairie, MN, USA) at a constant crosshead speed of 0.9 mm/min. The tensile specimens were machined along the rolling direction, and a minimum of three valid tests were performed for each material condition. EBSD data were processed and analyzed using the CHANNEL 5.0 software (Oxford Instruments, Abingdon, UK).

3. Results and Discussion

3.1. Accumulative Roll Bonding (ARB) Process

Figure 3 shows the microstructure of the composite after 6 ARB cycles. As seen in Figure 3a, no significant microstructural differences are observed between the Al12Zn2.2Mg1.7Cu layer and the Al12Zn2.2Mg1.7Cu3TiB2 layer, and the interfaces between the constituent layers are well-bonded [20]. The constituent layers can be distinguished by the presence of TiB2 particles, which are visible as fine clusters elongated along the rolling direction (RD). Bright white phases, identified as AlZnMgCu phases based on existing studies, are distributed within both layers. Figure 3b,c present TEM images of the Al12Zn2.2Mg1.7Cu layer and the Al12Zn2.2Mg1.7Cu3TiB2 layer, respectively. Dislocation walls consisting of parallel dislocations are evident in the microstructure of both constituents. In the Al12Zn2.2Mg1.7Cu3TiB2 layer, a limited number of dislocation intersections and a certain density of particles are also observed. In the ARB-processed sheet observed along the rolling direction-normal direction (RD-ND) plane, the interlayer interfaces remain essentially straight and continuous. Severe plastic deformation processes often induce varying degrees of interfacial plastic instability in laminated composites. However, in this study, the interfacial morphology of both composites showed no significant change with increasing ARB cycles. This is primarily attributed to the relatively high rolling temperature and the frequent intermediate annealing during the ARB process, which effectively suppressed the onset of plastic instability. This observation is consistent with the findings of L. F. Zeng et al. [18] for Cu/Ta laminates, who also reported that performing ARB at elevated temperatures effectively alleviates plastic instability.
The critical strain for the plastic instability of the harder layer during the rolling of multilayer composites can be predicted theoretically, based on factors such as the strain hardening exponent, strength coefficient, and the thickness of the constituent layers. To this end, Reihanian et al. [20] developed a theoretical model to predict the onset of plastic instability during the ARB process. In this study, their model was modified to better suit the current alloy system. The critical equivalent strain, ε*neck, for necking in the hard layer can be calculated using the following formula:
ε n e c k ∗ = n c 3 ( 1 − H ) 1 − A s A c ε n e c k ∗ n s − n c
In the formula, H is the ratio of the central constituent layer thickness to the total thickness of the composite sheet after N ARB cycles; n and A represent the strain hardening exponent and the strength coefficient of the constituent layers, respectively, where the subscripts c and s denote the central and surface constituent layers.
The basic flow properties of the two alloys used in this study are listed in Table 2. Based on the aforementioned formula and the relevant alloy data, the critical strain for the onset of plastic instability in the laminated composite was calculated to be 0.59. However, the microstructure observed in Figure 3 after 6 ARB cycles (corresponding to a cumulative equivalent strain of 3) shows no significant evidence of plastic instability. A comparison reveals that the theoretical prediction is considerably lower than the experimental observation. This discrepancy is primarily attributed to the intermediate annealing treatments applied during the ARB process. The high-temperature annealing effectively suppressed the tendency for necking at relatively low strains, thereby ensuring the successful progression of the ARB cycles.
The average thickness and hardness of the two constituents in the laminated composite are primarily dependent on the number of ARB cycles, in addition to their intrinsic physical properties. As shown in Figure 4, the thicknesses of both constituent layers exhibit a similar trend with increasing ARB cycles. After 6 cycles, the thickness of the Al12Zn2.2Mg1.7Cu layer is slightly lower than that of the Al12Zn2.2Mg1.7Cu3TiB2 layer. The thickness reduction per cycle is most pronounced during the initial ARB cycles and gradually levels off with further processing. Regarding hardness evolution, the hardness of the Al12Zn2.2Mg1.7Cu layer shows a slight decrease in the initial stage before stabilizing at a relatively constant level. In contrast, the hardness of the Al12Zn2.2Mg1.7Cu3TiB2 layer remains largely consistent throughout the ARB process. Typically, severe plastic deformation leads to significant strengthening in 7 series aluminum alloys due to strain hardening and grain refinement [21,22]. However, in this study, the hardness of the composite did not increase markedly with increasing strain. This is primarily attributed to the softening effect induced by the high rolling temperature and the repeated intermediate annealing. The annealing process promotes a dynamic balance between dislocation generation and annihilation, thereby counteracting the deformation strengthening effect [23]. Furthermore, for the Al12Zn2.2Mg1.7Cu layer in its initial state, recrystallization occurring during the high-temperature homogenization significantly alters the crystal structure, also contributing to the reduction in hardness.
Based on the experimental observations and analysis presented above, it is demonstrated that laminated composite sheets with similar mechanical properties and microstructures can be successfully fabricated using the ARB process combined with appropriate intermediate annealing. The results indicate that even after up to 6 ARB cycles, the properties of the constituent layers remain stable. Moreover, the interfacial bonding zones exhibit no signs of plastic instability, with the interface morphology maintaining a straight and continuous structure.

3.2. Cold Rolling Process at Room Temperature

3.2.1. Microstructure After Cold Rolling Process

Figure 5 displays the local morphology on the RD-ND plane of the 96-layer laminated composite (after 6 ARB cycles) under different cold rolling reductions. As shown in Figure 5a, after 48% cold reduction, the laminated interface between the two materials remains essentially straight. Figure 4b presents the corresponding EDS elemental mapping for Ti, indicating that the TiB2 particles are predominantly distributed within the Al12Zn2.2Mg1.7Cu3TiB2 layer. This further confirms the straight nature of the interface, which is primarily attributed to the relatively similar initial hardness of the two materials, resulting in minor differences in their flow stress and thickness ratio at low reductions [11,12,20]. With an increase in the cold rolling reduction, the interface gradually becomes unstable. At a reduction of 64%, the composite interface begins to exhibit plastic instability, characterized by slight necking in the Al12Zn2.2Mg1.7Cu layer; however, the interfacial continuity remains largely intact. As illustrated in Figure 5e, when the cold rolling reduction reaches 80%, the interfacial waviness becomes more pronounced. The thickness of the Al12Zn2.2Mg1.7Cu3TiB2 layer varies irregularly along its length, with localized regions showing significant necking. Notably, distinct lens-shaped fragment structures were not observed [10,24,25].
Figure 6 shows the average hardness and thickness values of the Al layer and the TiB2 layer after different cold rolling reductions. As seen in Figure 6a,b, the hardness of both the Al and TiB2 layers exhibits a gradual increasing trend with higher cold rolling reduction, and the trends for the two constituent layers are largely consistent. This can be attributed to the significant strengthening effect of strain hardening and grain refinement in metallic materials following severe plastic deformation. However, the presence of dynamic recovery during cold rolling leads to a dynamic balance between dislocation generation and annihilation within the laminated composite. Consequently, the mechanical properties of the deformed constituent layers do not show a substantial improvement. Furthermore, the thickness evolution reveals that the thinning of the Al layer in the normal direction (ND) is more pronounced than that of the TiB2 layer as the cold rolling reduction increases. Nevertheless, the thickness difference between the two layers remains relatively stable.
Figure 7 shows TEM micrographs of the constituent layers in the Al/TiB2 composite sheet after 80% cold rolling reduction. As shown in Figure 7a, the interaction and entanglement of dislocations under high strain conditions lead to an increase in dislocation density in both the Al12Zn2.2Mg1.7Cu layer and the Al12Zn2.2Mg1.7Cu3TiB2 layer. Additionally, grains with low dislocation density surrounded by those with high dislocation density are observed in both constituents, indicating the occurrence of dynamic recrystallization during processing [25]. Meanwhile, a large number of uniformly distributed plate-shaped second-phase particles are observed in both materials. EDS analysis indicates that these particles are primarily dispersed phases rich in Al, Zn, Cu, and Mg (AlZnMgCu phases) [26]. The formation of these second-phase particles strongly impedes dislocation motion during plastic deformation, thereby influencing the mechanical properties of the material. In the Al12Zn2.2Mg1.7Cu3TiB2 layer, the high strain induced by the cold rolling process results in a high density of dislocations around the TiB2 particles. The resulting strain fields and interactions between dislocations (indicated by yellow arrows) lead to the formation of fine, grayish-black banded-contrast images (BCI), as highlighted by the yellow circle in Figure 7b [27]. Analysis of the TEM microstructures of the two constituent layers reveals that the inherent structural differences between the original materials are a key factor contributing to the plastic instability observed in the composite sheet. Furthermore, based on the experimental observations, systematically increasing the strain level allows for the precise identification of the critical point for the onset of plastic instability. This approach is beneficial for subsequent investigations into the influence of this phenomenon on the mechanical properties of the laminated composite.
Figure 8a,e present EBSD micrographs of the laminated composite after cold rolling reductions of 16% and 80%, respectively. In these images, black and red lines represent high-angle grain boundaries (misorientation > 15°) and low-angle grain boundaries (misorientation 2–15°), respectively. A fibrous, deformed microstructure elongated along the rolling direction (RD) is clearly visible in both Figure 8a,e. This alignment becomes more pronounced with increasing deformation. A significant number of fine, equiaxed recrystallized grains are also observed. These recrystallized grains consist primarily of long, chain-like discontinuous dynamic recrystallization (DDRX) grains with similar crystallographic orientations [26,28]. The formation of this DDRX structure during cold rolling contributes positively to texture strengthening within the laminated composite, thereby enhancing the work hardening rate. Furthermore, in the Al-12Zn-2.2Mg-1.7Cu-3TiB2 layer, the DDRX grain structure near TiB2 particles is subjected to higher rolling strain, resulting in an increased dislocation density around these particles, which also positively influences the mechanical properties of the laminate [26]. In contrast, grains adjacent along the normal direction (ND) exhibit diverse orientations. Figure 8b,f show the corresponding recrystallization maps. Comparison reveals a significant increase in the fraction of recrystallized structures with greater deformation. This is primarily because the progressive increase in cold rolling reduction generates more dislocations and shear bands, providing abundant nucleation sites for recrystallization. The sharply rising dislocation density supplies a substantial driving force for recrystallization. The high stacking fault energy, combined with thermo-mechanical coupling effects, promotes dynamic recovery, leading to the formation of sub-grain structures that facilitate recrystallization nucleation and growth [29,30].
Figure 8c,d,g,h present the grain size distribution and grain boundary misorientation distribution for the composite after cold rolling reductions of 16% and 80%, respectively. The average grain size decreases significantly with increasing cold rolling reduction. At a reduction of 80%, 50% of the grains have diameters below 400 nm, nearly all grains are smaller than 1000 nm, and the average grain diameter is 419 ± 1 nm. Furthermore, the proportion of high-angle grain boundaries (HAGBs) increases with greater deformation. At the 80% reduction level, the HAGB fraction reaches 77.79%. Notably, the distribution of boundary misorientations exhibits a bimodal character, with a pronounced separation between low-angle and high-angle grain boundaries. Specifically, 40% of the boundaries have misorientation angles greater than 45°. Both the grain size and misorientation distributions confirm the formation of an ultrafine-grained structure after an 80% cold rolling reduction.
Figure 9 shows the EBSD micrograph on the TD-ND plane and the corresponding grain size distribution after 80% cold rolling reduction. As observed in Figure 9a, a substantial number of continuous dynamic recrystallization (CDRX) grains (indicated by white arrows) have formed not only within the Al-12Zn-2.2Mg-1.7Cu-3TiB2 layer and the Al-12Zn-2.2Mg-1.7Cu layer, but also at the interface [26,31]. This microstructural evolution can be explained as follows: Upon initial cold deformation, dislocation multiplication occurs extensively. The high stacking fault energy of the material facilitates dislocation cross-slip and climb, leading to the reorganization of dislocations into dislocation walls rather than random tangles. With continued cold rolling, dislocations glide to and are absorbed by these cell walls, causing the boundaries to become sharper and the dislocation density within them to increase, thereby transforming them into sub-grain boundaries. Concurrently, the dislocation density inside the sub-grains decreases significantly. Under the sustained cold rolling strain, adjacent sub-grains undergo relative lattice rotation to accommodate the deformation, which gradually increases the misorientation between them. Once the misorientation reaches a critical value (e.g., >10°~15°), the boundary between two adjacent sub-grains may become unstable and migrate, leading to their coalescence into a larger sub-grain with higher misorientation. A new CDRX grain is considered to have formed when a sub-grain boundary transforms into a high-angle grain boundary through this continuous process of rotation and coalescence [32]. Figure 9b displays the grain size distribution on the TD-ND plane after the 80% reduction. The average grain size measured on this plane is 442.5 nm, which is consistent with the result from the RD-ND plane. This further corroborates the formation of an ultrafine-grained structure after an 80% cold rolling reduction.

3.2.2. Effect of Cold Rolling Process on Mechanical Properties

Figure 10 shows the engineering stress–strain curves of the laminated composite under different cold rolling reductions. As observed, both the ultimate tensile strength (UTS) and yield strength (YS) initially increase and then decrease with increasing deformation. The sample with a 64% reduction achieves the maximum UTS of 670.5 MPa. For reductions up to 64%, the enhancement in mechanical properties is primarily attributed to two factors inherent in severe plastic deformation: strain hardening and grain refinement [10,11]. These two strengthening mechanisms are dominant in SPD-processed materials. At lower reductions, strain hardening serves as the primary strengthening mechanism. As the reduction increases, grain refinement strengthening gradually becomes more prominent [13]. Furthermore, the interfacial structure of the laminated composite significantly influences the strengthening outcome. Typically, cracks tend to propagate along weakly bonded interlayer interfaces, leading to reduced strength. The improved bonding strength achieved after accumulative roll bonding (ARB) is a key contributor to the increased tensile strength [33]. Moreover, in the Al/Al laminate, the continuous layered structure acts as a barrier to dislocation motion, resulting in a boundary strengthening effect. When the cold rolling reduction reaches 80%, the tensile strength declines, primarily due to the onset of necking (as shown in Figure 5e). The tensile properties of the multilayer composite can be theoretically predicted using a rule of mixtures approach:
σ = VAlσAl + VTiB2σTiB2
where VAl, VTiB2, σAl and σTiB2 represent the volume fractions and tensile strengths of the Al12Zn2.2Mg1.7Cu and Al12Zn2.2Mg1.7Cu3TiB2 constituent layers within the composite sheet, respectively. According to the rule of mixtures, the applied load in the laminated composite is primarily borne by the Al12Zn2.2Mg1.7Cu layer. However, when the cold rolling reduction reaches 80%, necking occurs in both constituent layers. This leads to a significant reduction in the load-bearing capacity of each layer during plastic deformation, which is the primary reason for the decline in both strength and ductility observed at the 80% reduction level.
Figure 10b indicates that the elongation of the samples gradually decreases with increasing cold rolling reduction. The variation in elongation is typically governed by the strain hardening behavior during cold rolling. In the laminated composite, increasing the cold rolling reduction leads to the continuous accumulation of dislocations and a concomitant rise in internal stress. This elevated internal stress promotes the initiation and propagation of cracks, resulting in reduced ductility. Apart from strain hardening, the interfacial structure is another critical factor influencing the composite’s ductility. Generally, the interfacial bonding strength increases with greater deformation. However, while a strong interface is beneficial for strength, it can detrimentally affect ductility. The well-bonded interfaces, particularly after significant reduction, may act as sites for crack nucleation or provide a direct path for crack propagation, thereby compromising the overall plastic deformation capability [34].

3.3. Fracture Analysis

Figure 11 shows the fracture surface morphology of the Al12Zn2.2Mg1.7Cu/Al12Zn2.2Mg1.7Cu3TiB2 composite sheet after a 16% cold rolling reduction. As shown in Figure 11a, the interfaces in the fracture surface remain straight and continuous at this reduction level. The fracture morphologies of the Al12Zn2.2Mg1.7Cu layer and the Al12Zn2.2Mg1.7Cu-3TiB2 layer exhibit distinct differences. A relatively high density of dimples is present in the Al12Zn2.2Mg1.7Cu layer, while fewer dimples are observed in the Al12Zn2.2Mg1.7Cu3TiB2 layer. Instead, large cleavage facets are evident adjacent to the TiB2 particles in the latter, as indicated by the yellow ellipse in Figure 11b. These fracture characteristics suggest a mixed ductile-brittle fracture mode. Furthermore, the presence of shear stresses at the interface during plastic deformation resulted in slight drag lines. A minor interlocking phenomenon is observed between the two constituents. Consequently, the dimples in the Al12Zn2.2Mg1.7Cu layer near the interface appear elongated or oriented towards the interface. The difference in the flow properties between the constituent metals is identified as the primary reason for the development of this in-plane shear stress at the interface during plastic deformation.
Figure 12 shows the tensile fracture surface of the laminated composite after an 80% cold rolling reduction. An uneven thickness distribution between the constituent layers is clearly observed in Figure 12a. Due to the occurrence of plastic instability, shear features are evident within both metal layers (Figure 12b). The mechanism for this phenomenon is as follows: during tensile testing, the Al12Zn2.2Mg1.7Cu3TiB2 layer, due to its inherently lower ductility, cracks initially and bears a significant portion of the load [23]. As the thickness of both constituent layers decreases within the plastic instability zones, their load-bearing capacity is considerably reduced. Consequently, these thinned, unstable regions are more prone to fracture, ultimately impairing the ultimate tensile strength of the composite. Figure 12c displays a high-magnification image of the fracture morphology. Second-phase particles can be observed at the bottom of some dimples. These particles are likely formed by interdiffusion of elements near the interface, as reported in previous studies. Furthermore, compared to the sample with a 16% reduction (Figure 11), the dimple size in the 80% reduction sample is significantly reduced. This observation is consistent with the high-strength, low-ductility characteristics typical of laminated composites subjected to severe deformation. The reduction in dimple size is primarily attributed to the strain hardening and grain refinement that occur during the intensive plastic deformation process [35].
Figure 13 illustrates the microstructural evolution during the ARB cycles and the subsequent cold rolling process. The interfaces remain straight throughout the ARB stage. With increasing cold rolling reduction, necking appears at the interfaces of the composite layers, which is induced by the flow property mismatch between the constituents. At low reduction levels, the internal shear stress generated at the interface between adjacent layers promotes the initiation of necking or the formation of shear bands in both the Al12Zn2.2Mg1.7C3TiB2 and Al-12Zn2.2Mg1.7Cu layers. As the cold deformation increases further, the internal shear stress significantly exceeds the yield strength of the metallic layers. This leads to the fracture of the necked regions within both constituents. At the maximum cold rolling reduction (Figure 13c), the size and morphology of the fragments are primarily determined by the spacing of the intersecting shear bands. This implies that if the length of the fractured Al12Zn2.2Mg1.7Cu3TiB2 layer segments exceeds this shear band spacing, further strain would, in theory, cause additional fragmentation of these segments. This process would exacerbate the plastic instability, ultimately adversely affecting the mechanical properties of the laminated composite.

4. Conclusions

The interfacial evolution during the accumulative roll bonding (ARB) process and its influence on the mechanical properties of Al12Zn2.2Mg1.7Cu/Al12Zn2.2Mg1.7Cu3TiB2 multilayer sheets were investigated. To achieve high-strength laminated composites, the material was subsequently cold-rolled to different thickness reductions. The main conclusions of this study are summarized as follows.
  • The Al/TiB2 laminated composite sheet was successfully fabricated via the accumulative roll bonding (ARB) process for 6 cycles at 460 °C. The elevated rolling temperature resulted in straight and continuous interlayer interfaces.
  • Varying degrees of plastic instability were observed in the Al12Zn2.2Mg1.7Cu3TiB2 layer during cold rolling. Necking initiated at various locations within the composite after a 64% reduction, resulting in wavy interfaces. Furthermore, severe necking approaching fracture occurred in the Al12Zn2.2Mg1.7Cu3TiB2 layer at an 80% cold rolling reduction.
  • The grain size progressively decreased with increasing cold rolling reduction. At an 80% reduction, 50% of the grains had diameters below 400 nm, and the average grain size was refined to 418.9 nm. Furthermore, the fraction of high-angle grain boundaries (HAGBs) reached 77.79%, with 40% of boundaries exhibiting misorientation angles greater than 45°. The grain size and misorientation distribution confirm the formation of an ultrafine-grained structure at the 80% reduction level, successfully yielding a nanocomposite sheet.
  • The strength of the sheet increased significantly with the cold rolling reduction, primarily due to strain hardening and grain refinement. However, a substantial decrease in strength occurred at the 80% reduction, which is attributed to the onset of severe necking.

Author Contributions

W.S.: Conceptualization, Writing—original draft; Z.X.: Methodology, Investigation; J.L.: Data curation, Methodology; Z.Y.: Data curation, Software; Y.H.: Software; Z.C.: Supervision, Writing—review and editing. All authors have read and agreed to the published version of the manuscript.

Funding

This work was supported by the Natural Science Foundation of China (No. 51871006).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Not applicable.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

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Materials 19 01031 g001
Figure 1. The experimental setup of the two-roll mill: (a) the feed inlet; (b) the side view of the rolling mill.
Figure 1. The experimental setup of the two-roll mill: (a) the feed inlet; (b) the side view of the rolling mill.
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Materials 19 01031 g002
Figure 2. Schematic illustration of the accumulative roll bonding (ARB) process for fabricating metallic laminated composites.
Figure 2. Schematic illustration of the accumulative roll bonding (ARB) process for fabricating metallic laminated composites.
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Materials 19 01031 g003
Figure 3. (a) The microstructure of the composite after 6 ARB cycles; (b,c) The TEM images of the two component layers.
Figure 3. (a) The microstructure of the composite after 6 ARB cycles; (b,c) The TEM images of the two component layers.
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Materials 19 01031 g004
Figure 4. The thickness and hardness changes in Al layer and TiB2 layer in different cycles of ARB process. (a) Al layer; (b) TiB2 layer.
Figure 4. The thickness and hardness changes in Al layer and TiB2 layer in different cycles of ARB process. (a) Al layer; (b) TiB2 layer.
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Materials 19 01031 g005
Figure 5. SEM image of RD-ND plane of laminated composite sheet after cold rolling deformation (a) 48% reduction; (c) 64% reduction; (e) 80% reduction; (b) The energy spectrum corresponding to (a); (d) The energy spectrum corresponding to the (c); (f) The energy spectrum corresponding to (e).
Figure 5. SEM image of RD-ND plane of laminated composite sheet after cold rolling deformation (a) 48% reduction; (c) 64% reduction; (e) 80% reduction; (b) The energy spectrum corresponding to (a); (d) The energy spectrum corresponding to the (c); (f) The energy spectrum corresponding to (e).
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Materials 19 01031 g006
Figure 6. The thickness and hardness of Al layer and TiB2 layer change with cold rolling deformation. (a) Al layer; (b) TiB2 layer.
Figure 6. The thickness and hardness of Al layer and TiB2 layer change with cold rolling deformation. (a) Al layer; (b) TiB2 layer.
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Materials 19 01031 g007
Figure 7. High-rate electron microscope images of the metal component layer after 80% cold rolling deformation: (a) Al; (b) TiB2.
Figure 7. High-rate electron microscope images of the metal component layer after 80% cold rolling deformation: (a) Al; (b) TiB2.
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Materials 19 01031 g008
Figure 8. EBSD microstructure of RD-ND plane of laminated composites, (a,e) IPF pictures; (b,f) recrystallization microstructure diagram; (c,g) statistics of grain size; (d,h) statistics of grain boundary angle, (Pale purple is low-angle grain boundaries, and pink represents high-angle grain boundaries).
Figure 8. EBSD microstructure of RD-ND plane of laminated composites, (a,e) IPF pictures; (b,f) recrystallization microstructure diagram; (c,g) statistics of grain size; (d,h) statistics of grain boundary angle, (Pale purple is low-angle grain boundaries, and pink represents high-angle grain boundaries).
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Materials 19 01031 g009
Figure 9. (a) EBSD microstructure of TD-ND plane of layered composites; (b) statistics of grain size.
Figure 9. (a) EBSD microstructure of TD-ND plane of layered composites; (b) statistics of grain size.
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Materials 19 01031 g010
Figure 10. (a) The engineering stress–strain curves of layered composites with different cold rolling deformations, (b) the variation trend of layered composites with cold rolling deformations.
Figure 10. (a) The engineering stress–strain curves of layered composites with different cold rolling deformations, (b) the variation trend of layered composites with cold rolling deformations.
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Materials 19 01031 g011
Figure 11. Scanning electron micrograph of the fracture surface of the laminated composite material after cold rolling deformation of 16%: (a) low magnification; (b) high times.
Figure 11. Scanning electron micrograph of the fracture surface of the laminated composite material after cold rolling deformation of 16%: (a) low magnification; (b) high times.
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Materials 19 01031 g012
Figure 12. SEM photos of fracture surface of laminated composites after 80% cold rolling deformation. (a) low multiple; (b,c) multiples.
Figure 12. SEM photos of fracture surface of laminated composites after 80% cold rolling deformation. (a) low multiple; (b,c) multiples.
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Materials 19 01031 g013
Figure 13. The development trend of the interface of layered composites with the increase in strain., (a) ARBed samples; (b) Cold rolling samples with deformation less than 64%; (c) Cold rolling samples with more than 64% deformation.
Figure 13. The development trend of the interface of layered composites with the increase in strain., (a) ARBed samples; (b) Cold rolling samples with deformation less than 64%; (c) Cold rolling samples with more than 64% deformation.
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Table 1. Chemical composition elements of TiB2/AlZnMgCu composites (wt. %).
Table 1. Chemical composition elements of TiB2/AlZnMgCu composites (wt. %).
AlloyZnMgCuZrTiBAl
A12.292.721.750.11200Bal.
B12.352.691.680.1211.610.82Bal.
Table 2. Flow properties of the constituents used in this study.
Table 2. Flow properties of the constituents used in this study.
MaterialWork Hardening
Exponent		Strength
Coefficient (MPa)		Tensile Strength (MPa)Tensile Strain for Fracture (%)
Al12Zn2.2Mg1.7Cu0.0819877438.5
Al12Zn2.2Mg1.7Cu3TiB20.0449147592.9
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MDPI and ACS Style

Sun, W.; Xiang, Z.; Li, J.; Yang, Z.; Han, Y.; Chen, Z. Effects of Cold Rolling on the Microstructure and Properties of Al/TiB2 Laminated Composites Fabricated by Accumulative Roll Bonding. Materials 2026, 19, 1031. https://doi.org/10.3390/ma19051031

AMA Style

Sun W, Xiang Z, Li J, Yang Z, Han Y, Chen Z. Effects of Cold Rolling on the Microstructure and Properties of Al/TiB2 Laminated Composites Fabricated by Accumulative Roll Bonding. Materials. 2026; 19(5):1031. https://doi.org/10.3390/ma19051031

Chicago/Turabian Style

Sun, Wenchao, Zhilei Xiang, Jihao Li, Zian Yang, Yang Han, and Ziyong Chen. 2026. "Effects of Cold Rolling on the Microstructure and Properties of Al/TiB2 Laminated Composites Fabricated by Accumulative Roll Bonding" Materials 19, no. 5: 1031. https://doi.org/10.3390/ma19051031

APA Style

Sun, W., Xiang, Z., Li, J., Yang, Z., Han, Y., & Chen, Z. (2026). Effects of Cold Rolling on the Microstructure and Properties of Al/TiB2 Laminated Composites Fabricated by Accumulative Roll Bonding. Materials, 19(5), 1031. https://doi.org/10.3390/ma19051031

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