X65 bimetallic sheets fabricated by explosive welding and hot rolling

    Microstructure and mechanical properties of CP-Ti/X65 bimetallic sheets fabricated by explosive welding and hot rolling Miao-Xia Xie, Lin-Jie Zhang, Gui-Feng Zhang, Jian-Xun Zhang, ZongYue Bi, Ping-Cang Li PII: DOI: Reference:

S0264-1275(15)30270-7 doi: 10.1016/j.matdes.2015.08.021 JMADE 429

To appear in: Received date: Revised date: Accepted date:

10 March 2015 29 July 2015 4 August 2015

Please cite this article as: Miao-Xia Xie, Lin-Jie Zhang, Gui-Feng Zhang, Jian-Xun Zhang, Zong-Yue Bi, Ping-Cang Li, Microstructure and mechanical properties of CPTi/X65 bimetallic sheets fabricated by explosive welding and hot rolling, (2015), doi: 10.1016/j.matdes.2015.08.021

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ACCEPTED MANUSCRIPT Microstructure and mechanical properties of CP-Ti/X65 bimetallic sheets fabricated by explosive welding and hot rolling Miao-Xia Xie a, Lin-Jie Zhang b, Gui-Feng Zhang b, Jian-Xun Zhang b,

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Zong-Yue Bi c, Ping-Cang Li d School of Mechanical and Electrical Engineering, Xi'an University of Architecture and Technology, Xi’an, 710055, China

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State key laboratory of mechanical behavior for materials, Xi’an Jiaotong University, Xi’an, 710049, China

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Baoji Petroleum Steel Pipe Co., Ltd., Baoji, Shaanxi, 721008, China

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Xi’an Tianli Clad Metal Materials Co., Ltd., Xi’an, 710201, China

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Abstract CP-Ti/X65-pipe-steel bimetallic sheet was fabricated by explosive welding and hot rolling (W&R). Trace of the wavy CP-Ti/X65 interface formed from explosive welding was observed along the straight CP-Ti/X65 interface of bimetallic sheet fabricated by W&R. The microstructure and component analysis showed the following. (i) The cross-section of the X65 zone consisted of a 2-5 µm-wide Ti diffusion layer next to the interface, a 150-200 µm-wide decarbonization layer, and the rest area with a banded structure morphology. (ii) There were numerous voids and a slight C element enrichment at the interface. (iii) The section of the CP-Ti zone consisted of a 10-50 μm-wide Fe diffusion zone next to the interface, a residual adiabatic shear band zone next to Fe diffusion zone, and the rest region composed of the α-Ti microstructure. The micro-hardness profile across CP-Ti/X65 interface was measured. The variation patterns of the mechanical properties of the bimetallic sheet in the thickness direction were obtained from stratified tensile tests. The shear test proved that the CP-Ti/X65 bimetallic sheet produced by W&R had acceptable shear bond strength. The microstructure and alloy element distribution across the TA1/X65 interfaces of as-welded, heat treated and extruded TA1/X65 bimetallic sheets were studied and compared. Keywords: CP-Ti/X65 bimetallic sheet; Explosive welding; Hot rolling; Microstructure; Mechanical property. 1. Introduction With the increasing reduction in oil and gas resources that contain relatively few impurity elements, there have been more and more demands from the petrochemical industry for the exploitation, long-distance transportation, and refinement of highly corrosive crude oil and natural gas. Due to their good corrosion resistance, high specific strengths, and good bearing capacities at elevated temperatures, Ti and its alloys have important applications and have already been widely used in industrial areas, such as aviation, aerospace, ships, and oil and petrochemical engineering. 

Corresponding author. Tel.: 86-29-82663115; Fax: 86-29-82663115 E-mail: [email protected] (L-J Zhang), [email protected] (M-X Xie)

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However, pure Ti and Ti alloys are far more costly than the commonly used low alloy steel or stainless steel. Therefore, structures that are made entirely from Ti and Ti alloys are extremely expensive. Commonly used low alloy steel often costs less than 2% of the price of Ti. Bimetallic sheets with steel as the base layer and commercially pure (CP)-Ti or Ti as the clad layer can significantly improve the service life of transportation pipelines and refining equipment and drastically reduce manufacturing costs. Explosive welding is one of the most widely used methods to manufacture bimetallic sheets. Currently, researchers have successfully performed explosive welding for hundreds of different dissimilar metal combinations, such as Ti/Ni[1, 2], Ni/Al [3], Cu/Ti [4] , Mg/Al [5], Ti/Al [6], W/CuCrZr [7], Cu/steel [8, 9], Inconel-alloy/steel [10], Al/steel [11], Zr/steel [12], stainless-steel/steel[13, 14] metallic-glass/steel [15, 16]. In recent years, in-depth, theoretical investigation of explosive welding process itself is quickly increasing, and the application of explosive welding technology in industry field is rising [17-26]. Hard, brittle Fe-Ti intermetallic compounds will be formed at the weld seams when traditional welding methods are used to weld Ti and steel. Explosive welding can form a metallurgical bonding interface between two metals; in addition, explosive welding will not result in significant crystallization and phase transition; therefore, explosive welding is extremely suitable for bonding Ti and steel. The microstructures, properties, and failure mechanisms of Ti/steel composite sheets fabricated using explosive welding are also attracting more and more attention. Nishid et al studied the mechanism responsible for bonding explosively welded Ti/steel clads [17]. The authors investigated the microstructural modifications of the bonding interface in explosively welded Ti/SUS 304 and Ti/SS 41 steel clads by transmission electron microscopy (TEM). Metastable phases, such as amorphous and supersaturated solid solutions of -Ti, were observed at the bonding interface in the explosively welded Ti/SUS 304 and Ti/SS 41 steel clads, which were considered as traces of melting and subsequent rapid solidification of thin layers along the contact surface of both parent materials. It was concluded that the melting layer was responsible for the bonding of the explosively welded Ti/steel clads. A. Chiba et al reported that the TiC layer was formed at the bonding interface of an explosively welded titanium/high-carbon steel plate after post-annealing, which served as a barrier for the diffusion of species across the interface and suppressed the formation of Fe-Ti intermetallic compounds [18]. As a result, a high bonding strength was preserved even after prolonged annealing at elevated temperatures. Y. Yang et al studied the microstructures of the adiabatic shear band (ASB) on the pure titanium side in a titanium/mild steel explosive welding interface by means of TEM [19]. It was found that extremely fine equiaxed grains (< 0.1 μm) with a low dislocation density were formed in the ASB, which were the result of dynamic recrystallization. The authors suggested that dramatic microstructural refinement by dynamic recrystallization enables a thermomechanical response that may lead to superplastic deformation in the ASB, which results in a high shear strain 2

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in the ASB. J. Song et al claimed that the interface region of explosively cladded Ti-steel joints consisted of four successive hierarchical levels [20]. It was concluded that the type, brittleness, and abutting microstructure of the intermetallic zones played the dominant role for the overall mechanical performance of the compounds when exposed to tensile and bending loads. In the work by P. Manikandan et al, commercially pure titanium and stainless steel were explosively welded by employing a thin interlayer of stainless steel [21]. The use of an interlayer was thought to reduce the kinetic energy loss and formation of the intermetallic layer at the interface. In the study by U. Kamachi Mudali et al, dissimilar joints produced by explosive joining and friction joining were compared and evaluated for applications in the severely corrosive nitric acid conditions used in reprocessing plants [22]. The results showed that direct titanium-304L SS joints manufactured using the explosive joining process exhibited good comprehensive properties. S.A.A. Akbari Mousavi presented an analytical calculation to determine the weldability domain or welding window [23]. The analytical calculations were in good agreement with experimental results. The welding conditions were tailored through the parallel geometry route with different explosive loads. Based on the element metallurgical compatibility, Chu et al designed Cu/V-based filler metals to join explosion-welded CP-Ti/Q345 bimetallic sheets [34]. Compared with the case using Cu/V solid wires, the number of crack defects could be reduced when a Cu-V flux-cored wire was used. This result was attributed to continuously distributed microstructures in the weld joint and low heat input with the Cu-V flux-cored wire. S.A.A. Akbari Mousavi et al found that post-heating of an explosively welded CP-Ti/AISI 304 composite within a temperature range of 650-950 ℃ formed different intermetallic phases at the joint interface [35]. Moreover, post-heating increased the width of the interfacial layers of the composite. Dariusz Rozumek et al studied fatigue crack growth in a Ti/steel composite under oscillatory bending. It was observed that fatigue crack growth occurred parallel to the applied loading, and its direction changed at the interface line. It was also confirmed that transcrystalline cracks were dominant at the fractures in the steel/titanium composite [36]. A. Karolczuk et al studied fatigue phenomena in explosively welded Ti/steel bimetallic plates subjected to push-pull loading [37]. It was found that fatigue crack initiation began at 80% in steel and at 20% at the interface (with no crack initiation in titanium). In addition, specimens with a flat interface exhibited a unique higher fatigue life compared to specimens with a wavy interface under a similar force amplitude. N. Kahraman reported that increasing the explosive ratio increased the deformation and hardness of both plates and therefore increased corrosion [38]. It was also found that the corrosion rate was high at the beginning of the experiment, and then, the rate of the corrosion decreased gradually during the experiment. N. Kahraman et al noted that the reason for the mass increment of the joined plates with the increase in explosive loads was that cold deformation increased in both metal surfaces from the explosion [39]. The brittle intermetallic results in lower strength of 3

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bonded joints. Therefore, reducing the volume fraction of brittle intermetallic is one of the key factors to improve the bonding strength. Researchers found that when a soft interlayer of Cu was used the formation of Fe-Ti intermetallic could be prevented and the bonding quality could be improved [40-44]. Copper interlayer can block the elements diffusion between titanium alloy and steel. Alhazaa et al joined two dissimilar aerospace alloys Ti-6Al-4V and Al7075 by using a combination of Cu coatings and Sn-3.6Ag-1Cu interlayers [45]. The results show that the Sn-3.6Ag-1Cu interlayer produces sound joints when a thin coating of Cu was deposited onto the alloy surfaces. Boiko et al showed that the technology of copper coating on steel and niobium coating on titanium in diffusion welding of titanium alloy to stainless steel is prospective [46]. In Zhao et al.’s study, the hot-roll bonding was carried out in vacuum between titanium alloy and stainless steel using niobium interlayer [47]. The results show that the plasticity of bonded joint is improved significantly. X65 pipe steel is a high strength, high toughness steel produced by thermo-mechanical controlled rolling. Using CP-Ti/X65-pipe-steel bimetallic sheet to construct long-distance, high-pressure oil or gas transportation pipelines can not only reduce the overall project cost due to the thinner wall thicknesses but also greatly extend the lifetime of transportation pipelines due to the high corrosion resistance of CP-Ti. In order to achieve CP-Ti/X65-pipe-steel bimetallic sheet with a large area and a thin clad layer, explosive welding and hot rolling compound process was adopted to fabricate bimetallic in the current work. The microstructure and mechanical properties of the bimetallic sheets achieved were studied. The TA1/X65 bimetallic sheets studied in this work would be used to produce long oil and gas pipelines and therefore would subjected to multiple bending process (i.e. JCOE process). This study focused on the inhomogeneity of both microstructure and mechanical properties of the obtained TA1/X65 bimetallic sheet, which would lead to highly heterogeneous deformation and strain rate of bimetallic sheet during JCOE forming process and increase the risk of local failure. With the detailed information about the mechanical properties of the bimetallic sheet, a more realistic numerical model of the JCOE processing of TA1/X65 bimetallic sheet will be established in the future and employed to carry out a simulation-based optimization of JCOE processing parameters. 2. Materials and methods 2.1 Materials and explosive welding process The chemical compositions of the materials used in the present work are shown in Table 1 and Table 2. The flyer plate and base plate were made from TA1 and X65 pipe steel, respectively. Flyer plate with a thickness of 8 mm and base plate with a thickness of 58 mm were explosively welded using an ammonium nitrate fuel-oil mixture as the explosive material. The final thickness of the explosively bonded CP-Ti/X65 bimetallic sheets was approximately 66 mm. Then, the explosively bonded CP-Ti/X65 bimetallic sheets were subjected to multi-pass hot rolling under air atmosphere to reduce the thickness. Prior to rolling, the work pieces were heated to 4

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8505 ˚C under air atmosphere and held at this temperature for 2 h. The temperature of work piece was 8505˚C when the rolling process began and was approximately 6105˚C when the rolling process ended. Descending distance of roller in each single-pass rolling process did not exceed 20% of the sheet thickness. CP-Ti/X65 bimetallic sheets with dimensions of 10,000×2,000×16 mm were eventually obtained. The final thickness of the flyer plate and base plate was approximately 2 mm and 14 mm, respectively.

Table 1 Chemical composition of X65 pipe steel (wt. %) C

Si

Cr

Mn

Ni

Cu

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0.053

0.33

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1.18

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0.14

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Table 2 Chemical composition of TA1 (wt. %) Fe

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Ti

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2.2 Metallographic studies Specimens for microstructural examination of the CP-Ti/X65 bimetallic sheet were prepared by standard metallographic method. The flyer plate was etched with an etchant consisting of 10 mL of HF, 10 mL of HNO3, and 80 mL of H2O, whereas the base plate was etched with an etchant consisting of 4 mL of HNO3 and 96 mL of alcohol. The samples were then examined under a Nikon Eclipse MA200 type optical microscope for general microstructural features and under a LS-JLLH-22 scanning electron microscope (SEM) to study the microstructural changes near the bond interface. Energy dispersive X-ray spectroscopy (EDS) analysis was performed to study the composition variation across the CP-Ti/X65 interface. 2.3 Mechanical test The micro-hardness of the CP-Ti/X65 bimetallic sheet was measured using a load of 300 g-force and a holding time of 15 s. Shear specimens were cut from the CP-Ti/X65 bimetallic sheet. Shear strength tests were performed on a universal mechanical testing machine at a loading speed of 0.2 mm/min. Charpy V-notch Impact Tests were conducted at -10˚C. To understand the changes in the mechanical properties of the CP-Ti/X65 bimetallic sheets in the sheet thickness direction, seven slices of the metallic material with a limited thickness of approximately 1.2 mm were cut from the CP-Ti/X65 bimetallic sheet in the horizontal direction, and the tensile test was performed for each metal slice. 2.4 Fracture observation Fractography studies on broken specimens of stratified tensile tests were performed using a LS-JLLH-22 SEM. The fracture characteristics of the CP-Ti/X65 bimetallic sheet were discussed based on SEM observation and EDS analysis of the distribution of elements on the fracture surface. 3. Results and discussion 5

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3.1 Optical microstructure As mentioned, the total thickness of the CP-Ti/X65 sheet was approximately 16 mm, and the TA1 in the clad layer had a thickness of approximately 2 mm. Fig. 1 shows the optical microstructure near the CP-Ti/X65 interface on the cross-section of the bimetallic sheet. It can be clearly seen from Fig. 1a that the CP-Ti/X65 interface was approximately a straight interface. It was reported that shear strength of the joint interface was highly dependent on the mechanical locking of the flyer and base plates provided by the wavy pattern [48]. Norris [49] found that there is no appreciable change in the shear strength of the rolled specimens even after heat treatment since any change in wave amplitude is not expected when a rolled sample is subjected to heat treatment. In this work, the amplitude of the as clad wavy pattern would decrease greatly due to hot rolling reduction from 66 to 16 mm and the extent of mechanical locking provided might be influenced and hence results in lower shear strength compared to that of an as welded joint.

Fig. 1 Optical microscope images of the interface in the W&R CP-Ti/X65 bimetallic sheet. (a) Straight interface, (b) high-resolution image (HRI) of position B in Fig. 1a, (c) HRI of position C in Fig. 1b, (d) HRI of position D in Fig. 1b.

As shown in Fig. 1a and b, residual morphology of peninsula (RMP), adiabatic shear band (RMASB), and vortices (RMV) can be easily identified from the 6

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cross-section of the CP-Ti/X65 bimetallic sheet. It can be deduced that the bonding interface of the as-welded CP-Ti/X65 bimetallic sheet had a wavy morphology, whereas the morphology of the bond interface had changed from wavy to straight during hot rolling. Due to the periodic characteristic of the interface in as-welded bimetallic sheet, the morphology of the interface in W&R bimetallic sheet also exhibited periodic characteristics. According to Fig. 1a, the distance between two neighboring RMPs was approximately 3 mm. As shown in Fig. 1a and b, the CP-Ti/X65 interface can be classified into two types based on the width of the transition zone near the interface: the majority of the CP-Ti/X65 interface exhibited a sharp transition, and the length of this type of interface accounted for approximately two-thirds of the total length of the interface; the transition zones between the two materials were relatively wide (approximately 50-100 µm) near locations where the RMPs were observed, and the length of this type of interface with a relatively wide transition zone accounted for approximately one-third of the total length of the interface. The residual morphology of the vortices accompanied by several voids can be clearly seen on the X65 side near the locations where the RMPs were observed, as shown in Fig. 1b. It can also be seen from Fig. 1b that the residual morphology of the ASB can be observed in the section of the TA1 zone that is close to the CP-Ti/X65 interface. A definition of adiabatic shear banding given by Hargreaves and Werner in 1974 [50] is that the phenomenon is ‘‘a thermo-mechanical instability which is established when the rate of strengthening due to strain hardening and strain-rate hardening is overcome by a combination of the rate of geometrical softening, due to void nucleation and growth, and the rate of thermal softening due to the conversion of mechanical work into thermal energy.’’ Two main types of adiabatic shear bands have been identified [51, 52]. The first are transformed bands and the second are deformed bands. It was found that the tendency of metals to form transformed or deformed bands could be mapped as a function of their thermal diffusivity and critical strain and strain rate for localization [50]. Ti alloys are known to readily form ASBs and fracture at high strain rates such as ballistic impact due to their high dependence of flow stress on temperature, their low strain hardening rate and poor thermal conductivity [53-55]. It was reported that the ASBs are a special case of thermoplastic instability favored by high rates of compressive strain [56]. Therefore, the huge pressure and shear force of the explosion welding process on the TA1 side near the interface produced many thin strip morphology that extended out from the interface, which were commonly known as fly line structures or shear bands. Fig. 1c and d shows the morphology of the microstructure of X65 at position C near the interface and the microstructure of TA1 at position D near the interface, respectively. As shown in Fig. 1c, the microstructure of the X65 included ferrite grains (brighter) and pearlite grains (darker) with a band pattern characteristic. It can be seen from Fig. 1c that the closer the section of the X65 zone was to the interface, the smaller the percentage of pearlite microstructure in this section, which might be 7

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due to the decarbonization of the X65 steel near the interface. Ti is more active than Fe and can easily seize C in the steel. Momono et al performed diffusion welding on CP-Ti and ultra-low carbon steel with a carbon content of only approximately 0.01 mass% at 900˚C; the results showed that there was a TiC layer at the interface of the joint [57]. The X65 pipe steel considered in this work had a carbon content of approximately 0.053 mass%. The as-welded bimetallic sheets were held at 8505˚C for 2 h after explosive welding, and then subjected to multi-pass rolling at an elevated temperature, which was beneficial for the diffusion of C from X65 to CP-Ti. In addition, intense plastic deformation occurred in the metals near the interface of the bimetallic sheet, and a large amount of lattice defects could be formed inside the metals during the explosive welding process and the subsequent hot-rolling process, which provided diffusion channels for C. It can be seen from Fig. 1d that a relatively significant recovery and recrystallization occurred in the section of the TA1 zone near the interface (upper part of Fig. 1d), whereas the section of the TA1 zone relatively far from the interface was less affected, which might be because there were relatively more lattice defects in the section of the TA1 zone near the interface after explosive welding; in other words, the final grain size of titanium resulted from not only the temperature of the work piece but also the density of lattice defect [35]. From the above analysis, it can be seen that the morphology of the microstructure near the CP-Ti/X65 interface was relatively complex. Residual morphologies of peninsulas, islands, vortices, and voids near the interface would significantly affect the subsequent bimetallic-pipe manufacturing processes, such as plate bending, plate rolling, and diameter extension. 3.2 SEM observation and composition analysis across the CP-Ti/X65 interface Three representative positions near the CP-Ti/X65 interface (indicated by B, C, and D) are shown in Fig. 2a. The transition zone between the two materials near position B was extremely narrow. The transition zone between the two materials near position C was relatively wide, and RMPs existed near position C. The transition zone between the two materials near position D was relatively wide, and the residual morphologies of the islands existed near position D. Fig. 2b, c and d shows the microstructure near positions B, C, and D, respectively.

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Fig. 2 SEM images of the interface in the CP-Ti/X65 bimetallic sheet. (a) Typical morphology of

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the CP-Ti/X65 interface, (b) high-resolution image (HRI) of position B in Fig. 2a, (c) HRI of

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position C in Fig. 2a, (d) HRI of position D in Fig. 2a, (e) HRI of position E in Fig. 2a.

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Fig. 3a shows the backscattered electron (BSE) image of the section near the CP-Ti/X65 interface. Fig. 3b, c, and d show the line scan analysis results of the three typical positions. Fig. 3e, f, and g shows the micro-zone analysis results of the three typical positions. The probe size of the EDS analysis was approximately 4 μm in diameter. According to Fig. 2 and Fig. 3, the zone near the interface can be classified into four layers (indicated by I, II, III, and IV, respectively, in Fig. 3a) in the direction perpendicular to the CP-Ti/X65 interface. The X65 zone above the CP-Ti/X65 interface is area I; Area III is characterized by the residual morphology of the ASB; Area II is located between the CP-Ti/X65 interface and the ASB zone; and Area IV is the section of the CP-Ti zone where a certain recovery and recrystallization occurred. It can be seen from Fig. 3a that the brightness of Area II was lower than that of Area I but higher than those of Area III and Area IV.

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Fig. 3 Distribution of elements across the CP-Ti/X65 interface. (a) BSE image, (b) distribution of alloy elements (DAE) at position B in Fig. 3a, (c) DAE at position C in Fig. 3a, (d) DAE at position D in Fig. 3a, (e) micro-zone analysis results (MAR) of position B in Fig. 3a, (f) MAR of position C in Fig. 3a, (g) MAR of position D in Fig. 3a.

Fig. 2b, Fig. 3b and e show the test results of the section near position B. As mentioned in the previous subsection, Area I was composed of ferrite (F) and pearlite (P) and exhibited a banded microstructure morphology. The characteristics of the banded microstructures were weakened in a 100-200 µm-wide zone near the interface, as shown in Fig. 2a and b. Fig. 3e clearly shows that there is a Ti diffusion layer in X65 with a width of approximately 3-5 µm near the interface. It can be seen from Fig. 2b and Fig. 3b that there were a small number of voids at the CP-Ti/X65 interface near position B. From Fig. 3b and e it was found that the Fe diffusion layer in TA1 was located in the section of Area II near position B; this diffusion layer near position B had a width of approximately 5-10 µm, and Fe had a mass percentage of greater than 8% in the diffusion zone. Considering that Fe is a β-phase stable element, it is 10

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believed, based on the analysis, that the section of Area II near position B might be composed of the supersaturated solid solution formed from the high-temperature β-phase microstructures that were preserved until the temperature returned to room temperature [35, 57-60]. The residual morphology of the ASB can be observed in the section of Area III near position B, as shown in Fig. 2b and Fig. 3b. Since that strict conditions are required for the formation of ASB, the morphology of ASB in Fig. 2b and Fig. 3b can only be formed during the explosive welding process. Existing research shows that Area III is composed of α-Ti microstructures [19]. As discussed in previous subsection, Area IV was the section of the α-Ti zone where a certain recovery and recrystallization occurred [35, 57]. Fig. 2c, Fig. 3c and f show the test results of the section near position C. It can be seen from Fig. 2 and Fig. 3 that the morphology of the microstructure in Area I, III, and IV near position C is similar to that near position B. The section of Area II near position C had a width greater than 50 µm, as shown in Fig. 2c. From Fig. 3f, it was found that the section of Area II near the X65 side had a Fe content of approximately 22%, while the section of Area II near the ASB side had a Fe content of approximately 1%. It can be seen from Fig. 2c that the section of Area II near X65 exhibited a smooth morphology under SEM, whereas the section of Area II close to the ASB side exhibited acicular microstructure morphology. Similar microstructure morphologies (i.e. smooth morphology near steel side and acicular morphology near Ti side) have been found by other researchers near the interface in both explosive welded and diffusion bonded Ti/steel joint [35, 57]. Analysis shows that the stabilizing effect of Fe on the β phase allowed the high-temperature β-phase microstructures in the section of Area II near the X65 side to be preserved until the temperature returned to room temperature to form a supersaturated β-solid solution [35, 57]. The section of Area II near the ASB side had a relatively low Fe content, which was not sufficiently high to allow the high-temperature β-phase microstructures to be entirely preserved until the temperature returned to room temperature, and thus, the morphologies of α+β microstructures eventually formed [35, 57-60]. In other words, the bright β-Ti was embedded in the shaded α-Ti matrix. It can also be seen from Fig. 3f that there appears to be a slight C enrichment phenomenon at the interface (i.e. target (8) in Fig. 3f). Although the tabulated content of C in Fig. 3f was qualitative rather than quantitative, this might indicate that TiC was formed at the CP-Ti/X65 interface [18, 57]. In addition, it can be seen from Fig. 2a that there were significantly more voids at the CP-Ti/X65 interface near position C than at the CP-Ti/X65 interface near position B, which might reduce the bonding strength of these two materials. Fig. 2d, Fig. 3d and g show the test results of the section near position D. It can be seen from Fig. 2 and Fig. 3 that the morphology of the microstructure in Area II, III, and IV near position D is similar to that near position C. Figs. 3a and d show that the Ti contents were relatively high near the islands, which indicates that these islands 11

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were formed from a part of the TA1 material that became detached from the TA1 clad layer and entered the X65 base layer during the explosive welding process. The wavy interface formation in explosive welding resulted from the combined action of the detonation force and the metal vortex flow. As mentioned above, TA1 has a high dependence of flow stress on temperature, a low strain hardening rate, a low conductivity, and a low tensile strength. Therefore, when the detonation force and the metal vortex flow were very intense, a part of the material might separate from the flyer plate and enter the base metal and form an island-like morphology. It can also be seen from Fig. 3g that O was detected at the island morphology, which might be because the surface of the work piece was not completely cleaned or the air captured during the explosive welding process contained oxygen. An interesting phenomenon revealed by Figs. 1-3 is that a relatively significant recovery and recrystallization occurred in the section of the TA1 zone far from the interface (area below the ASB zone in Fig. 1d), while the residual morphology of ASB zone which was relatively close to the interface still be observed after the hot rolling process. The heating temperature and dislocation density are important factors that affect recovery and recrystallization. Yang et al investigated the microstructures of the adiabatic shear band on the pure titanium side in a titanium/mild steel explosive welding interface using TEM. Extremely small equiaxed grains (< 0.1 μm) of α-Ti with a low dislocation density in the grain boundary were observed in the ASB. Yang et al confirmed that the α-Ti -Ti transformation did not occur in the ASB, and the material in the ASB did not melt [19]. The low dislocation density in the ASB zone might be the reason why the morphology of the ASB could still be maintained after the hot rolling process. High-density dislocation will generally form in materials near an explosive welding interface, and such a high-density dislocation will provide channels for element diffusion. An ASB zone with a low dislocation density near such an explosive welding interface is clearly good for prohibiting the unfavorable diffusion process. Actually, it can be seen from Figs. 3b, c, and d that the ASB zone seems to have indeed prevented Fe from diffusing to TA1. If ASB can be proven to have a preventative effect on diffusion, then ASB near the interface might be beneficial for preventing the formation and growth of hard, brittle intermetallic compound layers near the CP-Ti/X65 interface and improving the bonding strength of the joint. From the above analysis, It can be found that the morphologies of the microstructures and the element distribution in the area near the interface on the TA1 side were much more complex than those on the X65 side, which was primarily related to the significant differences between the thermophysical properties and mechanical properties of the two materials [19]. Due to its relatively low strength, significant deformation could easily occur in TA1 during the explosive welding process, which would result in the formation of ASB. Due to the relatively poor thermal conductivity of TA1, local melting zones (LMZ) and heat-affected zone 12

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(HAZ) could be easily formed on the Ti side. In addition, Ti is relatively active and has a hexagonal close-packed lattice structure, which is favorable for the diffusion of Fe atoms in Ti, resulting in that diffusion distance of Ti in X65 was significantly smaller than that of Fe in TA1 [35]. These phenomena would significantly influence the subsequent pipe manufacturing processes of the bimetallic sheet, such as bending, plate rolling and tube diameter expanding. 3.3 Mechanical properties 3.3.1 Micro-hardness

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Fig. 4 Micro-hardness profile near the interface.

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Figs. 4a and b show the test locations and results of the Vickers hardness near the interface. It can be seen from Fig. 4a that the section of the X65 zone far from the interface had a Vickers hardness value of approximately 190 HV, and the section of the TA1 zone far from the interface had a Vickers hardness value of approximately 130 HV. However, micro-hardness values of more than 300 HV were measured at the locations where the RMP was observed, which was significantly greater than the micro-hardness values of the section of the X65 zone far from the interface and the section of the TA1 zone far from the interface. It can be seen from Fig. 4b that the test points with high hardness values were located in Area II (shown in Fig. 1). According to the previous discussion, Area II was composed of the supersaturated solid solution or hard, brittle intermetallic compounds. Such a high-hardness layer near the interface would result in the incompatibility of the strain at the interface under an external load, which in turn would result in a decrease in the bearing capacity of the joint. During hot rolling process, plastic flow of the metals would be generated near the interface under intense pressure. Because of the mechanical inhomogeneity across TA1/X65 interface, voids would be produced by separation of metal flow around hard and brittle intermetallic compounds [50]. 3.3.2 Stratified tensile test

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Fig. 5 Stratification scheme of the stratified tensile test for the W&R CP-Ti/X65 bimetallic sheet.

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Fig. 6 Dimensions of a stratified tensile test specimen.

Fig. 7 Stress-displacement curves of the W&R CP-Ti/X65 bimetallic sheet obtained from the

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Fig. 8 Variations in the mechanical properties of the W&R CP-Ti/X65 bimetallic sheet in the sheet thickness direction.

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Fig. 9 Morphologies of the specimens before and after the stratified tensile test.

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A stratified tensile test was performed on the W&R CP-Ti/X65 bimetallic sheet. Fig. 5 shows the stratification scheme of the sample collection. Tensile test specimens were fabricated according to the dimensions shown in Fig. 6. Fig. 7 shows the stress-displacement curves obtained from the tensile test. Fig. 8 summarizes the variation patterns of the tensile strength and the elongation percentage of the bimetallic sheet in the sheet thickness direction. Fig. 9 shows the morphologies of the specimens before and after the tensile test. The test results show that the material in layer 1 of the W&R CP-Ti/X65 bimetallic sheet had the lowest strength and the best plasticity. The material in layer 1 (i.e., TA1) had a tensile strength of approximately 330 MPa and an elongation percentage of approximately 40%; the material in layer 2 (i.e., CP-Ti/X65) had a tensile strength of approximately 500 MPa and an elongation percentage of approximately 22%; the material in layer 3 (i.e., X65) had a tensile strength of approximately 660 MPa and an elongation percentage of approximately 26%; and there was no significant difference in the mechanical properties between the material in layer 3 and the materials in layers 4, 5, and 6. The material in layer 7, which was close to the lower surface of X65, had a tensile strength of approximately 620 MPa and an elongation percentage of only 22%. 3.3.3 Shear test Shear test was performed on the CP-Ti/X65 bimetallic sheet. Fig. 10a shows the dimensions of shear test specimen. Fig. 10b is a schematic diagram of the shear test. In order to discuss the joint interface strength, shear test was conducted under the conditions that the angle between shear direction and detonation direction was 0°, 45° and 90° respectively, and the corresponding morphologies and element distribution of the shear fractures were shown in Fig. 11. Shear test was repeated three times under each condition. Fig. 12 shows the typical load-displacement curves achieved from shear tests.

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Fig. 10 Schematic diagram of (a) the dimensions of the shear test specimen and (b) shear test.

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It can be seen from Fig. 11a, b and c that the morphology of all the sheared fractures obtained under different conditions exhibited periodic characteristics. Energy dispersive X-ray spectroscopy analysis was carried out along the yellow line segments marked in Fig. 11a, b and c, and the corresponding results were given in Fig. 11d, e and f, respectively. When the angle between shear direction and detonation direction was 0° or 45°, the chemical composition of the most areas of the fracture was in between TA1 and X65 steel, as shown in Fig. 11d and e, which indicated that the fractures were consisted of intermetallic compounds. When the angle between shear direction and detonation direction was 90°, a considerable part of the fracture was within the TA1 plate, as shown in Fig. 11f. Therefore, the bonding joint exhibited the highest shear strength when the angle between shear direction and detonation direction was 90°, as shown in Fig. 12.

Fig. 11 Morphologies and element distribution of the shear fractures achieved when the angle between shear direction and detonation direction was (a) 0°, (b) 45°and (c) 90°, respectively. 16

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Fig. 12 Load-displacement curve from the shear test.

Fig. 13 Schematic diagram of formation of the sheared fracture morphology when the angle between shear direction and detonation direction was 90°.

As mentioned, the morphology of the interface in W&R bimetallic sheet had periodic characteristics. For the CP-Ti/X65 interface where the transition zones between the two materials were relatively wide (approximately 50-100 µm) and many voids were observed, crack might initiated on the I/II interface and propagated in the X65 material of which the strength was lower that of the intermetallic compounds in 17

ACCEPTED MANUSCRIPT the Area II, as shown in Fig. 13c and d. For the majority of the CP-Ti/X65 interface where exhibited a sharp transition and good bonding quality, the crack might initiated in the TA1 side and propagated in the TA1 material of which the strength was lower than the neighbored X65 material, as shown in Fig. 13b and d.

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Charpy V-notch Impact Tests were conducted at -10˚C on a standard Charpy Impact Test machine. Fig. 14 shows the dimensions of a Charpy V-notch Impact Test specimen. Three specimens were tested and the average Charpy impact energy was about 221 J.

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Fig. 14 Dimensions of Charpy V-notch Impact Test specimen.

Fig. 15 Morphologies and element distribution of the peeled fractures achieved by Charpy V-notch Impact Test.

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Fig. 15a is the morphology of the specimen after impact test, from which it can be found that the flyer plate near the impact fracture was peeled off the base plate. The side view of the peeled fracture in TA1 side (i.e., indicated by the arrows in Fig. 15a) is shown in Fig. 15 b. The fracture given in Fig. 15b exhibited a smooth morphology, which is quite different from those observed in Fig. 11a, b and c. This phenomenon could be related to the difference in loading direction and loading rate between shear test and impact test. Figs. 15c and d show the high-magnification images of the morphologies of positions C and D that are marked in Fig. 15b. From Fig. 15b, c and d, it can be deduced that the chemical composition of most parts of fracture was in between TA1 and X65 steel, indicating that the peeled fracture was mainly consisted of intermetallic compounds. 3.4 Analysis of the fractures of the stratified tensile test specimens Fig. 16a shows the macroscopic morphology of the fracture of the material in layer 1 (i.e., flyer plate CP-Ti). Figs. 16b, c, d, and e show the high-magnification images of the morphologies of positions B, C, D, and E that are marked on the image of the macroscopic morphology of the fracture. As shown in Figs. 16b and d, a large amount of small-scale dimples and a small amount of large-scale dimples can be observed in the central location of the fracture. It can be seen from Figs. 16c and e that the morphologies of the shear dimples can be observed everywhere at positions C and E near the edges of the macroscopic fracture. The SEM observation results (Fig. 16) show that the material in layer 1 had extremely good toughness. On the one hand, TA1 has good inherent toughness; in addition, material of layer 1 was far from the CP-Ti/X65 interface and was thus not affected by diffusion. On the other hand, recovery and recrystallization occurred during the hot rolling process, which was also beneficial for improving the toughness. It can be seen from Fig. 17a that bending deformation curving toward the relatively low-strength TA1 occurred in the fracture of the material in layer 2. The CP-Ti/X65 interface was completely exfoliated. As mentioned in the previous subsection, there was a hard, brittle Fe diffusion zone near the interface on the TA1 side, and this zone was usually near to the RMP and had a width greater than 50 µm. Song et al detected the FeTi phase at the locations where the peninsula morphologies were observed at the interface of the explosively cladded CP-Ti/Steel joint and discovered a large amount of micro-cracks at the peninsula morphologies after explosive welding. They found that these micro-cracks provided the cracking origin when the sheet was under load. In addition, there was a significant difference in the mechanical properties between the large blocks of the hard, brittle FeTi phase and the surrounding matrix, which led to insufficient plastic compliance of the intermetallic inclusion relative to the surrounding matrix, which, in turn, could mean that the interface might begin to crack at the locations where the RMPs were observed and eventually became exfoliated [20]. In this work, the clear morphology of a cleavage fracture was also observed near the location of RMP, as shown in Figs. 17e and f. To 19

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improve the bonding strength of the interface of the composite sheet, it is evidently necessary to reduce or eliminate the locations where the RMPs were observed at the

Fig. 16 Fractography of layer 1 (CP-Ti) after the tensile test. (a) Macro-fractography of layer 1, (b) high-resolution fractography (HRF) of position B in Fig. 16a, (c) HRF of position C in Fig. 16a,

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(d) HRF of position D in Fig. 16b, (e) HRF of position E in Fig. 16a.

Fig. 17 Fractography of layer 2 ( CP-Ti/X65 ). (a) Macro-fractography of layer 2, (b) high-resolution fractography (HRF) of position B in Fig. 17a, (c) HRF of position C in Fig. 17a, 20

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CP-Ti/X65 interface as much as possible. As shown in Fig. 17h, most of the areas of the X65 fracture exhibited dimple-shaped morphologies; in addition, an area with a width of approximately 5-10 µm in the fracture of X65 close to the interface exhibited a smooth morphology, which might be related to the occurrence of Ti diffusion in this area. It can also be seen from Fig. 17h that the CP-Ti/X65 interface exhibited an uneven morphology after becoming exfoliated, and no large-sized dimples were observed, which is similar to that observed by Song et al [20]. It can be seen from Figs. 17b, c, and d that the fracture on the TA1 side primarily exhibited a dimple-shaped morphology.

Fig. 18 Element distribution on the fracture near the peninsula morphologies. (a) Results of line scanning on the fracture near the peninsula morphology, (b) high-resolution fractography (HRF) of position B in Fig. 18a, (c) HRF of position C in Fig. 18a, (d) HRF of position D in Fig. 18a, (e) HRF of position E in Fig. 18a, (f) target areas of the EDS analysis, (g) EDS analysis results (EAR) of target (3) in Fig. 18f, (h) EAR of target (2) in Fig. 18f, (i) EAR of the target (1) in Fig. 18f.

Line scanning analysis and micro-zone analysis were performed on the fracture near the locations where the RMPs were observed. Fig. 18 shows the analysis results. It can be seen from Fig. 18 that there were fragments with diameters of approximately 100 µm and cleavage fracture morphologies in the facture near the locations where the peninsula morphologies were observed; the analysis showed that these fragments were composed of approximately 52.33 mass% Fe and approximately 46.27 mass% Ti, 21

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indicating that local melting occurred and that solidification areas consisting of a mixture of these two metals were formed at locations where the peninsula morphologies were observed during the explosive welding process. It has been reported by many researchers that intermetallic compounds could be formed at the Ti/steel interface where the peninsula morphology was observed [17, 20, 35]. Based on the analysis of the Ti-Fe binary phase diagram, it is believed that intermetallic compound FeTi might be formed at the locations where the peninsula morphologies were observed after the peninsulas had solidified and cooled down. Hard, brittle FeTi may easily develop micro-cracks and may also easily result in strain stratification when bearing load, which, in turn, decreases the strength of the interface. It can be seen from Figs. 18b and e that the HAZs near the local melting area exhibited dimple-shaped morphologies, which might be because recovery and recrystallization occurred in the HAZs, and the existence of such HAZs might, to a certain degree, inhibit the cracks of the intermetallic compound area from propagating into the matrix [20]. As shown in Fig. 18a, there were significant fluctuations in the contents of the two elements on the X65 side, which was because plastic flow of the metals near the peninsula morphologies was generated under intense pressure, which could easily become vortical flow under the restraints of the surrounding wavy matrix material [20], resulting in a mixed, staggered distribution of the two materials.

Fig. 19 Fractography of layer 3 (X65). (a) Macro-fractography of layer 3, (b) high-resolution fractography (HRF) of position B in Fig. 19a, (c) HRF of position C in Fig. 19a, (d) HRF of position D in Fig. 19a, (e) HRF of position E in Fig. 19a, (f) HRF of position F in Fig. 19a, (g) HRF of position G in Fig. 19a.

Fig. 19a shows the macroscopic morphology of the fracture of material in layer 3. Figs. 19b, c, and d show the magnified images of positions B, C, and D that are marked in the macroscopic image of the fracture. It can be seen from Fig. 19b, c, and d that the areas inside the fracture far from the central part of the fracture all exhibited 22

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dimple-shaped microscopic morphologies. As shown in Fig. 19e, f, and g, there were a large amount of dimples in the central location of the fracture, and lamellar cracking and step-like fracture morphologies could also be observed. The analysis shows that because the material in layer 3 was relatively close to the TA1/X65 interface, the material in layer 3 underwent relatively significant hardening after cladding, which resulted in the formation of lamellar cracking and step-like fracture morphologies. The difference between the morphologies of the central location of the fracture and the areas on the two sides of the fracture was primarily attributed to the fact that the central location of the fracture and the areas on the two sides of the fracture were subjected to restraint conditions with different intensities during the tensile process, and the metals at the central location were under stricter restraint conditions. 3.5 Analysis of the effects of heat treatment and extrusion The hot rolling process has both thermal and mechanical effects on bimetallic sheet. In this subsection, in order to get a little bit more understanding about the effects of hot rolling process on CP-Ti/X65 bimetallic sheet, heat treatment test (85020 ˚C, 2h, under air atmosphere) and mechanical extrusion test (25 ˚C, 20% reduction in thickness) were conducted respectively on the explosively welded CP-Ti/X65 bimetallic sheet with a total thickness of about 31mm. The microstructural morphology of interfacial intermetallics and the DAE across TA1/X65 interface in as-welded, heat treated and extruded TA1/X65 bimetallic sheets were studied and compared. Figs. 20, 21 and 22 show the test results for as-welded, heat treated and extruded TA1/X65 bimetallic sheet, respectively.

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line L6 in Fig. 20c, (k) DAE along line L7 in Fig. 20c, (l) DAE along line L8 in Fig. 20c.

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As shown in Fig. 20b, the bonding interface of as-welded TA1/X65 bimetallic sheet had a wavy morphology. Fig. 20c, g-l show that a substantial part of the as-welded TA1/X65 interface can be identified by a sharp transition between the TA1 and the X65 materials and the region over which there was a significant change in Ti content was approximately 2-3 microns wide. Fig. 20d shows the transition zone between TA1 and local melted zone, from which it can be seen that the width of the transition zone between TA1 and the local melted zone was less than 5 microns. Namely, alloy element diffusion in as-welded TA1/X65 joint was quite limited.

Fig. 21 Test results for heat-treated TA1/X65 bimetallic sheet. (a) Macrograph of cross-section, (b) Wavy interface, (c) Typical morphology of the wave interface, (d) high-resolution image (HRI) and micro-zone analysis results (MAR) of position D in Fig. 21c, (e) Distribution of alloy elements (DAE) along line L1 in Fig. 21c, (f) DAE along line L2 in Fig. 21c, (g) DAE along line 24

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Fig. 21 shows the test results for heat-treated TA1/X65 bimetallic sheet. As shown in Fig. 21 b, e and f, the region over which there was a significant change in Ti content was approximately 100 microns wide after heat treatment, which is much larger than that observed in the as-welded TA1/X65 bimetallic sheet ( i.e. Fig. 20). It can be seen from Fig. 2 c and d that the section of transition zone near X65 exhibited a smooth morphology, whereas the section of transition zone close to the TA1 side exhibited acicular microstructure morphology. As shown in Fig. 21d, similar microstructure morphologies have also been found in heat-treated TA1/X65 bimetallic sheet. From this, formation of the acicular microstructure morphology in the transition zone of the W&R TA1/X65 bimetallic sheet should be attributed to the heat treatment process before rolling.

Fig. 22 Test results for extruded TA1/X65 bimetallic sheet. (a) SEM image of TA1/X65 interface, (b) Cracks at position B in Fig. 22a, (c) Micro voids at position C in Fig. 22b, (d) DAE at position D in Fig. 22a, (e) Micro voids at position E in Fig. 22d.

Fig. 22 shows the test results for extruded TA1/X65 bimetallic sheet. As shown in Fig. 22b, a number of cracks were observed at the CP-Ti/X65 interface of extruded bimetallic sheet. In Fig. 22c and e, it can be seen that there were many micro voids in the transition zone after mechanical extrusion and the TA1 zone far away from CP-Ti/X65 interface is free of voids. During hot rolling process of TA1/X65 25

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bimetallic, cracks and voids may also occur at those weak positions arrowed in Fig. 22b, c and e. As shown in Fig. 2b, c and d, there were a number of voids at the interface of W&R TA1/X65 bimetallic sheet. In summary, the TA1/X65 interface of an explosively welded bimetallic sheet exhibited reliable, robust shear strength. There was a significant difference in the mechanical properties between the materials in the clad layer and the base layer of the TA1/X65 bimetallic sheet. Residual morphologies of hard, brittle intermetallic compounds, hole defects, peninsulas, and islands existed near the interface. Such characteristic will significantly impact the subsequent JCOE pipe manufacturing process. It is necessary to adjust the parameters of explosive welding and reduce the formation of the peninsula morphologies and local melting and solidification areas near the interface after explosive welding as much as possible; Also, role of the adiabatic shear band near the interface of bimetallic sheet might be a topic worthy of in-depth study. 4. Conclusions In this work, the microstructure and mechanical properties of explosive welded TA1/X65 pipe steel bimetallic sheet which would be used to manufacture JCOE welded oil and gas pipe were investigated, with emphasis on the inhomogeneity of both microstructure and mechanical properties of the bimetallic sheet obtained. The following conclusions can be drawn from this study: 1. The TA1/X65 bimetallic sheet fabricated using the W&R method had a straight TA1/X65 interface; periodically occurring RMPs could be observed at the interface. 2. The morphology of the microstructure near the interface exhibited periodic characteristics in the direction parallel to the TA1/X65 interface. The interface can be classified into two types based on the morphological characteristics: the interface with an extremely narrow transition zone, whose accumulative length accounted for approximately two-thirds of the total length of the interface, and the interface with a relatively wide transition zone, which often occurred near the locations where the RMPs were observed and had an accumulative length that accounted for approximately one-third of the total length of the interface. 3. The area near the interface can be classified into four areas in the direction perpendicular to the TA1/X65 interface based on the morphological characteristics of the microstructures: (1) Area I, which is the section of the X65 zone that is above the TA1/X65 interface; (2) Area III, which exhibits the morphology of the residual ASB; (3) Area II, which is located between the TA1/X65 interface and the ASB; and (4) Area IV, which is the section of the TA1 zone where a certain recovery and recrystallization occurred. 4. The shear test results show that the composite sheet had acceptable shear strength. Variation patterns of the mechanical properties of the bimetallic sheet 26

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This work was supported by the National High Technology Research and Development Program of China (Grant No. 2013AA031303HZ). The authors would like to thank Ms. Jie Ning, Ms. An Wang, Mr. Qing-Lin Bai and Mr. Jian-Nan Yang from Xi’an Jiaotong University for their help in optical microscopy, SEM and mechanical testing.

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ACCEPTED MANUSCRIPT List of Figure Captions Fig. 1 Optical microscope images of the interface in the W&R CP-Ti/X65 bimetallic sheet. (a) Straight interface, (b) high-resolution image (HRI) of position B in Fig. 1a, (c) HRI of

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position C in Fig. 1b, (d) HRI of position D in Fig. 1b.

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Fig. 2 SEM images of the interface in the CP-Ti/X65 bimetallic sheet. (a) Typical morphology of the CP-Ti/X65 interface, (b) high-resolution image (HRI) of position B in Fig. 2a, (c) HRI

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of position C in Fig. 2a, (d) HRI of position D in Fig. 2a, (e) HRI of position E in Fig. 2a. Fig. 3 Distribution of elements across the CP-Ti/X65 interface. (a) BSE image, (b) distribution of alloy elements (DAE) at position B in Fig. 3a, (c) DAE at position C in Fig. 3a, (d) DAE at position D in Fig. 3a, (e) micro-zone analysis results (MAR) of position B in Fig. 3a, (f)

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MAR of position C in Fig. 3a, (g) MAR of position D in Fig. 3a. Fig. 4 Micro-hardness profile near the interface.

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Fig. 5 Stratification scheme of the stratified tensile test for the W&R CP-Ti/X65 bimetallic sheet. Fig. 6 Dimensions of a stratified tensile test specimen.

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stratified tensile test.

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Fig. 7 Stress-displacement curves of the W&R CP-Ti/X65 bimetallic sheet obtained from the

Fig. 8 Variations in the mechanical properties of the W&R CP-Ti/X65 bimetallic sheet in the sheet

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thickness direction.

Fig. 9 Morphologies of the specimens before and after the stratified tensile test. Fig. 10 Schematic diagram of (a) the dimensions of the shear test specimen and (b) shear test.

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Fig. 11 Morphologies and element distribution of the shear fractures achieved when the angle between shear direction and detonation direction was (a) 0°, (b) 45°and (c) 90°, respectively.

Fig. 12 Load-displacement curve from the shear test. Fig. 13 Schematic diagram of formation of the sheared fracture morphology when the angle between shear direction and detonation direction was 90°. Fig. 14 Dimensions of Charpy V-notch Impact Test specimen. Fig. 15 Morphologies and element distribution of the peeled fractures achieved by Charpy V-notch Impact Test. Fig. 16 Fractography of layer 1 (CP-Ti) after the tensile test. (a) Macro-fractography of layer 1, (b) high-resolution fractography (HRF) of position B in Fig. 16a, (c) HRF of position C in Fig. 16a, (d) HRF of position D in Fig. 16b, (e) HRF of position E in Fig. 16a. 31

ACCEPTED MANUSCRIPT Fig. 17 Fractography of layer 2 (CP-Ti/X65). (a) Macro-fractography of layer 2, (b) high-resolution fractography (HRF) of position B in Fig. 17a, (c) HRF of position C in Fig. 17a, (d) HRF of position D in Fig. 17a, (e) HRF of position E in Fig. 17a, (f) HRF of position F in Fig. 17a, (g) HRF of position G in Fig. 17a, (h) HRF of position H in Fig.

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Fig. 18 Element distribution on the fracture near the peninsula morphologies. (a) Results of line scanning on the fracture near the peninsula morphology, (b) high-resolution fractography

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(HRF) of position B in Fig. 18a, (c) HRF of position C in Fig. 18a, (d) HRF of position D in Fig. 18a, (e) HRF of position E in Fig. 18a, (f) target areas of the EDS analysis, (g) EDS analysis results (EAR) of target (3) in Fig. 18f, (h) EAR of target (2) in Fig. 18f, (i) EAR of the target (1) in Fig. 18f.

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Fig. 19 Fractography of layer 3 (X65). (a) Macro-fractography of layer 3, (b) high-resolution fractography (HRF) of position B in Fig. 19a, (c) HRF of position C in Fig. 19a, (d) HRF

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of position D in Fig. 19a, (e) HRF of position E in Fig. 19a, (f) HRF of position F in Fig. 19a, (g) HRF of position G in Fig. 19a.

Fig. 20 Test results for as-welded TA1/X65 bimetallic sheet. (a) Macrograph of cross-section, (b)

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Wavy interface, (c) Typical morphology of the wave interface, (d) high-resolution image (HRI) and micro-zone analysis results (MAR) of position D in Fig. 20c, (e) Distribution of

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alloy elements (DAE) along line L1 in Fig. 20c, (f) DAE along line L2 in Fig. 20c, (g) DAE along line L3 in Fig. 20c, (h) DAE along line L4 in Fig. 20c, (i) DAE along line L5 in Fig. 20c, (j) DAE along line L6 in Fig. 20c, (k) DAE along line L7 in Fig. 20c, (l) DAE along

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line L8 in Fig. 20c.

Fig. 21 Test results for heat-treated TA1/X65 bimetallic sheet. (a) Macrograph of cross-section, (b) Wavy interface, (c) Typical morphology of the wave interface, (d) high-resolution image

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(HRI) and micro-zone analysis results (MAR) of position D in Fig. 21c, (e) Distribution of alloy elements (DAE) along line L1 in Fig. 21c, (f) DAE along line L2 in Fig. 21c, (g) DAE along line L3 in Fig. 21c, (h) DAE along line L4 in Fig. 21c, (i) DAE along line L5 in Fig. 21c, (j) DAE along line L6 in Fig. 21c, (k) DAE along line L7 in Fig. 21c, (l) DAE along line L8 in Fig. 21c. Fig. 22 Test results for extruded TA1/X65 bimetallic sheet. (a) SEM image of TA1/X65 interface, (b) Cracks at position B in Fig. 22a, (c) Micro voids at position C in Fig. 22b, (d) DAE at position D in Fig. 22a, (e) Micro voids at position E in Fig. 22d.

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Graphical abstract

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Highlights Big-size TA1/X65 sheet is achieved by explosive welding & hot rolling

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Periodic characteristic of microstructure across TA1/X65 interface is studied

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Mechanical inhomogeneity of TA1/X65 bimetallic sheet is studied

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Variation of chemical composition across TA1/X65 interface is studied

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Proposal abound for improving performance of TA1/X65 bimetallic sheet is given

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