Butt laser welding-brazing of AZ31Mg alloy to Cu coated Ti-6Al-4V with AZ92 Mg based filler

Butt laser welding-brazing of AZ31Mg alloy to Cu coated Ti-6Al-4V with AZ92 Mg based filler

Optics and Laser Technology 117 (2019) 200–214 Contents lists available at ScienceDirect Optics and Laser Technology journal homepage: www.elsevier...

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Optics and Laser Technology 117 (2019) 200–214

Contents lists available at ScienceDirect

Optics and Laser Technology journal homepage: www.elsevier.com/locate/optlastec

Full length article

Butt laser welding-brazing of AZ31Mg alloy to Cu coated Ti-6Al-4V with AZ92 Mg based filler

T



Jinge Liua, Caiwang Tana,b, , Laijun Wub, Xiaoye Zhaoa,b, Zequn Zhanga, Bo Chenb, Xiaoguo Songa,b, Jicai Fenga a b

State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China Shandong Provincial Key Laboratory of Special Welding Technology, Harbin Institute of Technology at Weihai, Weihai 264209, China

H I GH L IG H T S

Mg and Cu-coated Ti using laser welding-brazing process. • Join of Cu coating thickness on welding quality and interfacial reaction was studied. • Influence elements diffusion mechanism on Mg/Cu-coated Ti interface was clarified. • AJoint • strength and fracture path and were associated with interfacial bonding and thickness.

A R T I C LE I N FO

A B S T R A C T

Keywords: Laser welding-brazing Titanium alloy Magnesium alloy Metallurgical bonding Interfacial reaction

Butt laser welding-brazing of Mg to Ti with Mg based filler was performed with the assistance of Cu coating. The thickness of Cu coating was varied in the study, to investigate its influence on microstructure and mechanical properties of the joint. Two main regions were distinguished along the Mg/Cu coated Ti interface. At the upper interface, the bonding mechanism evolved from mechanical bonding into a “Ti3Al and AlCu2Ti” metallurgical bonding, and into a “Ti3Al + Ti2Cu + AlCu2Ti” metallurgical bonding as Cu coating thickness varied from 0 μm to 19.7 μm to 24.9 μm. At the lower interface, mechanical bonding changed into the IMC layer consisted of Ti3Al, Ti2Cu and AlCu2Ti when Cu coating thickness reached 28.2 μm. The chemical potential calculated by the developed Toop model suggested that Cu played a role in promoting the mutual diffusion between Al and Ti. Ti tended to react with Al first compared with Cu suggesting Ti-Al compounds were produced easily at Mg/Ti interface. In the case of Cu coating thickness of 19.7 μm, the joint load reached a peak value of 3457 N, as high as 85.35% of that of Mg base metal. The fracture mode changed from interfacial failure to fusion zone failure when the coating thickness reached 19.7 μm.

1. Introduction Recently, a surge of interest has been paid to facilitate lightweight industrial fabrication in order to reduce the exhaust emission and fuel consumption [1–3]. AZ31B magnesium (Mg) alloys as one of the most widely used magnesium alloys has received much attention due to its compelling qualities including good formability, low density, high strength-weight ratio, good malleability [1]. Titanium (Ti) alloy has been an advanced alloy for military, biomedicine, and aerospace industry which was proved to have preeminent mechanical and physical characteristics such as good corrosion, abrasion resistance, and thermal resistance. Ti-6Al-4V as a dual-phase (α + β) alloy with aluminum and vanadium included was acknowledged as a widely used titanium alloys ⁎

[4–6]. Since lightweight industrial fabrication is an effective approach for energy saving and environmental conservation, the joining of lightweight alloys of AZ31B and Ti-6Al-4V will broaden application prospects. However, there is still a challenge for joining magnesium alloy to titanium alloy because of the enormous differences in the properties in physics and metallurgy [7]. The melting points of pure Mg and Ti are 922 K and 1941 K, respectively. In addition, Mg will suffer severe vaporization when direct joining with Ti by conventional fusion welding, because the boiling point of pure Mg (1364 K) is much lower than the melting point of Ti (1941 K). Furthermore, the maximum solid solubility of Mg in Ti is close to nil, and that of Ti in Mg is only 0.12 wt%. Hence, intermediate element should be employed to address the

Corresponding author at: Shandong Provincial Key Laboratory of Special Welding Technology, Harbin Institute of Technology at Weihai, Weihai 264209, China. E-mail address: [email protected] (C. Tan).

https://doi.org/10.1016/j.optlastec.2019.04.024 Received 12 January 2019; Received in revised form 25 March 2019; Accepted 14 April 2019 0030-3992/ © 2019 Elsevier Ltd. All rights reserved.

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used as shielding gas. Table 2 lists other main process parameters employed in the LWB process.

problem in the metallurgical bonding of Mg/Ti joint. Joining of dissimilar metals Mg to Ti was realized in previous studies by transient liquid phase (TLP) brazing [8–11], friction stir welding (FSW) [12,13], cold metal transfer (CMT) welding [14], tungsten inert gas (TIG) welding [15,16] and laser keyhole welding [17,18]. In the TLP bonding process [11], Ni and Cu sandwich interlayers were used to perform TLP brazing of AZ31 Mg alloy and Ti-6Al-4V alloy. CuMg2, Mg2Ni and Mg3AlNi2 were noticed at the Mg/Ti interface, confirming that the addition of Cu and Ni played a role in promoting metallurgical bonding of Mg/Ti. In addition, the thickness of interlayer was proved to affect microstructural evolution. The interfacial reaction became violent producing a thicker IMC layer with the increase of Cu and Ni interlayers. Then, Atieh and Khan [9] added Cu nanoparticles dispersing in the Ni coating during Mg/Ti joining. They found that Cu nanoparticles influenced solidification rate which led to the joint-strength enhancement. In the FSW process [12], TiAl3 interfacial compound was detected because of the reaction between Ti base metal and Al element diffused from Mg alloy base metal. The tensile testing indicated the growth of this interfacial layer was harmful to joint mechanical properties. Regarding the process of CMT welding [14], the Al element in AZ61 filler played a role in improving metallurgical bonding at the Mg/ Ti interface through reacting with Ti substrate. In the case of TIG welding [15], a solid solution zone was dispersed at the Mg/Ti interface, suggesting that element diffusion occurred without formation of any intermetallic compound. Besides, ultrasonic-assisted treatment was found to play a role in grain refinement [16], which enhanced the joint strength during the TIG welding process. With regard to laser welding [17], reliable joining of Mg/Ti was achieved through a precise control of laser offset during laser keyhole welding. The acceptable joint appeared when the laser beam was shifted to 0.2–0.3 mm close to Ti side, which reached the fracture load of 266 MPa. Laser welding-brazing (LWB) has advantages over conventional fusion welding in high flexibility, low residual stress, and thinner IMC layer especially when joining dissimilar materials, such as Al/steel [18–20], Al/Ti [21,22], and Mg/Ti [23]. According to existing study [24], Ni coating on titanium alloy base metal was observed to improve the wettability of molten filler significantly. Additionally, Ni-Mg-Al ternary compound formed at the Mg/Ti interface, suggesting that metallurgical bonding was realized with the assistance of Ni coating. However, as a kind of strategic resources, the output of nickel alloy every year is limited. Copper and its alloy as a common metal have a reserve of 700 million ton and an output about 20 million ton, which are over ten times than that of nickel (data in 2016). Additionally, from the existing study [11], the Cu element added as a form of sandwich foils had a tendency of bilateral diffusion. Based on these analyses, Cu was selected as another potential interlayer element to further control the welding cost and improve the interfacial reaction. The present work aimed at the microstructural and mechanical characteristics of laser welding-brazing AZ31B/Ti-6Al-4V with various Cu interlayer thickness. The joining mechanism of Mg to Ti via Cu-coating with AZ92 filler was elucidated according to above analysis.

2.2. Electrodeposition process Before welding, a three-minute acid pickling (15% HCl, 5% HF and 80% distilled water) for titanium alloy was employed to clear the surface oxides away. Electrolytic pure copper the were electroplated on titanium base metals straight after acid pickling. The schematic of the Cu electroplating process was similar with our previous research about the LWB process of Mg/Ni-coated Ti [23,24]. The process of electrodeposition was carried out in a 700-mL plastic beaker. The plating solution in beaker were prepared by dissolving 120 g K4O7P2, 10 g C6H5O7(NH4)3, 12 g HNa2O4P, 24 g CuSO4·5H2O in 500 mL distilled water. During the process of electrodeposition, the cleaned copper sheet was the anode and titanium base metal was the cathode. The current density used for electrodeposition was set at 1.0 A/dm2, which was controlled by a constant-current source. The plating solution was agitated by a multifunctional magnetic stirrer 350 revolutions per minute (RPM). As the increase of coating time (30 min, 60 min, 90 min, and 120 min), the coating thickness changed as the red line shown in Fig. 2(a). The SEM morphologies of Cu coated Ti substrate with varied coating thickness δ (10.8 μm, 19.7 μm, 24.9 μm, 28.2 μm) were shown in the Fig. 3(b)–(e), which suggested that the Cu coating on Ti surface was uniform. 2.3. Analysis methods After the LWB process, scanning electron microscopy (SEM) and optical microscopy (OM) were employed to observe the interfacial microstructures and cross sections of the welded-brazed specimens made by standard metallographic preparation procedures respectively. Energy-dispersive spectrometry (EDS) and X-ray diffraction (XRD) techniques were employed to analyze the phase composition and confirm the specific reaction products at the interface. As shown in Fig. 1(b), the tensile-shear test specimens were cut into a size of 100 mm × 10 mm. The tensile test was carried out at room temperature, with a crosshead speed of 0.5 mm/min. The final tensile result was calculated by at least three tensile specimens. 3. Numerical simulation and thermodynamic analysis 3.1. Numerical simulation To obtain the interfacial thermal history, finite element simulation software MSC.Marc was employed. To make a balance between the accuracy and speed of calculation, finer meshes with a size of 0.16 × 0.16 × 0.2 mm3 in molten filler and heat affected zone (HAZ) nearby while coarser meshes in other regions were divided as shown in Fig. 3. All the meshes consisted of hexahedron elements and eight nodes. The total number of elements and nodes were 241,500 and 297,000 for the joint with V-shape interface respectively. To simulate the process of actual filler-feeding, the “birth and death element” method was adopted. The temperature-dependent physical properties of Mg and Ti [25] base metals are shown in Fig. 4. The density of TC4 and AZ31B was regarded as constant at the value of 4.51 g/cm3 and 1.8 g/cm3 respectively. The physical properties of weld seam were assumed as the same as Mg base metal. Due to the large distance of focus position employed during the LWB process, Gauss plane heat source was employed on the surface of the weld seam to match the characteristics of welding heat process according to the existing research [26]. The Gauss plane heat model could be determined by

2. Experimental 2.1. Laser welding-brazing process A 6-kW fiber laser with a beam parameter product of 7.2 mm × mrad and a KUKA robot with six-axis was employed in LWB process. The focused laser beam had a diameter of 0.2 mm and a wavelength of 1070 nm. A 200-mm lens was used to focus 200-μm core diameter fiber for laser beam transmission. Commercially available Ti6Al-4V sheets with a groove of 45°and AZ31B sheets, both with a size of 100 × 50 × 1.5 mm3 were used as base metals. A 1.2-mm-diameter AZ92 filler was selected as filler wire. The chemical compositions of Ti6Al-4V, AZ31B and filler are shown in the Table 1. Fig. 1(a) shows the schematic diagram of LWB process in butt configuration, pure Ar was

qs (x , y ) = 201

3α (x 2 + y 2 ) 3αQs exp[− ] πrs2 rs2

(1)

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Table 1 Chemical compositions of base metals and filler metal (wt%).

AZ31B Ti-6Al-4V AZ92

Al

Zn

Mn

Fe

V

Si

Mg

Ti

2.5–3.5 5.5–6.8 8.3–9.7

0.5–1.5 3.5–4.5 1.7–2.3

0.2–0.5 – 0.15–0.5

< 0.005 0.3 < 0.005

– 3.5–4.5 –

0.1 – < 0.005

Bal. – Bal.

– Bal. –

where Qs and γs2 are the effective power and effective radius of Gauss plane heat source respectively, α is the heat efficiency of welding process. x and y are the coordinates in the Gauss plane. The welding direction was along the x axis, so x could be defined as “x = Vt”. Where V was the welding speed and t was the welding time.

Table 2 Welding parameters in present work.

3.2. Thermodynamic analysis To analyze the atomic diffusion, thermodynamic calculation was performed using MATLAB. Toop model [27] and Miedema model [28] were used to analyze the interfacial reaction. The binary system’s formation enthalpy could be expressed as:

ΔH1,2 = f1,2

f1,2 =

Experimental parameters

Value

Distance of focus position from Ti surface Laser power Velocity of welding Flow rate of shielding gas Ar Velocity of wire feeding

+20 mm 1500 W 0.5 m/min 15 L/min 2.5 m/min

GE =

−1

1/3 + (ΔnWS )2

−1

(3)



x4 G E (x , 1 − x 1 14 1

E + (x3 + x 4 )2G34 (x

Gm =

G ID

μi =

∂Gm ∂x i

+

GE

x3 3 + x4

,

1 − x1)

x4 ) x3 + x 4

(6)

(7)

G12E = ΔH1,2 [1 − T (1/ Tm, i + 1/ Tm, j )/14]

(8)

GijE

In the formula (5), is acquired based on the theory of Tanaka (1990). G ID is the Gibbs energy in ideal condition of solution approximation. 4. Results 4.1. Joint appearances and cross sections

x2 x3 ΔH12 (x1, 1 − x1) + ΔH13 (x1, 1 − x1) 1 − x1 1 − x1 x2 x3 ⎞ + (x2 + x3 )2ΔH23 ⎛ , ⎝ x2 + x3 x2 + x3 ⎠

1 − x1) +

(5)

where xi is the molar fraction of different elements; nws, V and φ are the electron density parameters, molar volume and electronegativity of component; ΔH is the formation enthalpy of binary alloys; p, q, r, a, μ are experimental constants according to the research of Miedema and coworkers [28]. The other physical parameters used in Miedema model was listed in Table 3. To further gain the thermodynamic calculation result of ternary system Toop model [27] developed form Miedema model was employed. The formation enthalpy was calculated through the formula as:

ΔH123 =

x3 G E (x , 1 − x 1 13 1

(2)

1/3 2 2pV12/3 V22/3 [q/ p (ΔnWS ) − (Δφ)2 − a (r / p)] 1/3 (ΔnWS )1

1 − x1) +

x x E + G23 ( x +2 x , x +3 x ) 2 3 2 3 x x E + (x2 + x 4 )2G24 ( x +2 x , x +4 x ) 2 4 2 4

x1 [1 + μ1 x2 (φ1 − φ2)] x2 [1 + μ 2 x1 (φ2 − φ1)] x1 V12/3 [1 + μ1 x2 (φ1 − φ2)] + x2 V22/3 [1 + μ 2 x1 (φ2 − φ1)]

x2 G E (x , 1 − x 1 12 1

The coating thickness had an influence on the wettability of molten filler on titanium base metal which affected the weld appearances of Mg/Ti joint [24]. Fig. 5 shows the appearances of the Mg/Cu coated Ti joints produced with different coating thickness. A narrow weld with small wetting-spreading area was observed on the front surface as shown in Fig. 5(a), suggesting a poor wettability of Mg on bare Ti and lower fracture load. While, with the addition of fewer Cu coating of 10.8 μm, Ti sheet was still not sufficiently wet by filler metal leading to excessive deposition of filler on front and back surfaces of the joint. In



(4)

As for quaternary system, a model developed from Toop model [27] was employed. μi was corresponding chemical potential and GE was corresponding excess Gibbs free energy which was calculated by following formula.

Fig. 1. Schematic of LWB process of Mg to Cu-coated Ti: (a) LWB process; (b) specimen for tensile-shear test. 202

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Fig. 2. Micrograph of Cu coating on Ti substrate versus coating times: (a) effect of electrodeposition time on coating thickness; (b) 30 min; (c) 60 min; (d) 90 min; (e) 120 min.

19.7 μm, 24.9 μm, 28.2 μm) are presented in Fig. 7. Table 4 lists the corresponding EDS results to confirm the phase components. The FZ-Ti interface was divided into two typical zones, named the upper interface and lower interface, to investigate the microstructural evolution of Mg/ Ti interface. Fig. 7(a)–(c) shows the morphologies with back-scattering SEM when the coating thickness was less than 10.8 μm. At both upper and lower interfaces, no obvious IMC layer was observed. A large number of voids were observed in the fusion zone close to the Mg/Ti interface, which was detrimental to the bond strength. In addition, as the EDS line result of Mg/Ti interface shown in Fig. 7(c), only a diffusion zone of Mg and Ti was observed at the interface. The Cu and Al were not enriched at the interface, which indicated that the metallurgical reaction did not happen at the Mg/Ti interface. The main bonding mechanism of the interface was a fragile mechanical bonding with low strength. The similar mechanical bonding mechanism was also discussed in our previous research about the LWB process of Mg/bare Ti [25,30]. Before taking part in interfacial metallurgical reaction, the Cu coating was so thin that all the molten Cu got involved into the molten filler with the action of vortex produced in the molten pool. According to the previous investigation in Mg/Ti laser keyhole welding for unequal thickness butt joint [16], the white dot phase (P1) was observed to disperse in α-Mg, which was inferred as Mg17Al12. Meanwhile, strip grey phase (P2) was identified as Mg17(Al,Cu)12. The formation of this Mg-Al-Cu ternary

addition, compared with the joint appearance of Mg and bare Ti, an apparently wider weld was observed at the front and back surfaces, which could be concluded that Cu coating improved the wettingspreading ability of weld seam. However, with the Cu coating thickness further over 24.9 μm, a trend of collapse was observed at the front surface of joint, as indicated in Fig. 6(d) and (e). Fig. 6 shows the representative cross sections of Mg/Cu coated Ti joints with different Cu coating thickness, which further explained the influence of coating thickness. From Fig. 7(a)–(e), with the coating thickness δ changed from 0 μm to 28.2 μm, the weld reinforcement presented a descending trend changing from 1.46 mm to 0.22 mm. Meanwhile, the corresponding contact angle on Ti sheet changed from 79.1° to 0°, suggesting that Cu coating played a role in promoting the wettability of molten filler and this promoting effect increased with the thicker coating thickness. The similar influence of coating layer, such as Ni and Zn, was reported in our previous studies [24,29]. 4.2. Interfacial microstructure During the LWB process, the characteristic of high thermal gradient gave rise to the variation of intermetallic compound along the Mg/Ti interface, and thus further analysis was required [29]. The SEM morphologies and the magnifications of Mg/Cu coated Ti joints at different positions produced with various Cu coating thickness (10.8 μm,

Fig. 3. Finite element analysis model of Mg/Ti butt joint. 203

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Fig. 4. Temperature-dependent physical properties of Mg and Ti: (a) Mg; (b) Ti.

process. The amount of strip grey phase was relatively few, due to the limited addition of Cu coating at the coating thickness of 10.8 μm. Fig. 7(d)–(f) shows the morphologies with SEM back-scattering when the Cu coating thickness was 19.7 μm. When Cu coating was thick enough, Cu element took part in interfacial reaction gradually. An obvious IMC layer appeared at the upper interface, which indicated that the bonding mechanism of the Mg/Cu coated Ti interface converted from mechanical bonding to metallurgical bonding. At the upper interface, the interfacial microstructure presented a form of sandwich foils divided into three regions for research as shown in Fig. 7(d): the inside IMC layer (Layer I), a transition layer (Layer II) close to FZ and the FZ. The microstructure with higher magnification of layer I was shown in Fig. 7(f). The massive grey phase (P4) mainly contained 5.97 at% Cu, 10.68 at% Al and 82.23 at% Ti, which was identified as Ti3Al. Meanwhile, the white phase (P5) containing 26.58 at% Cu,

Table 3 Element parameters used in Miedema and Toop model. Element

Tm/K

nws/d.u.

φ/V

μ

V/cm3

Cu Al Mg Ti

1356 933 922 1933

3.18 2.7 1.6 3.51

4.55 4.2 3.45 3.8

0.07 0.04 0.1 0.04

7.72 10 14 10.58

phase was mainly ascribed to that the Cu atoms involved in the molten pool replaced some of the Al atoms in second phase β-Mg17Al12 during cooling process. Similar phenomenon about the formation of Mg17(Al,Zn)12 was also reported by Gao et al. [17]. In their research, the Zn atoms could replace Al atoms in Mg17Al12 in precipitation

Fig. 5. Appearances of Mg/Cu coated Ti butt joints with various coating thicknesses: (a) 0 μm; (b) 10.8 μm; (c) 19.7 μm; (d) 24.9 μm; (e) 28.2 μm. 204

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Fig. 6. Cross sections of Mg/Cu coated Ti butt joints with various coating thicknesses: (a) 0 μm; (b) 10.8 μm; (c) 19.7 μm; (d) 24.9 μm; (e) 28.2 μm.

mapping results shown in Fig. 8. In addition, Al element was more enriched in IMC layer suggesting that the supplement of Cu may promote the metallurgical reaction of Al and Ti. Fig. 7(j)–(l) shows the morphologies with SEM back-scattering when the Cu coating thickness was 28.2 μm. As shown in Fig. 7(j) the upper interface of IMC/FZ was extremely uneven. In the IMC layer, snowflakelike phases were observed in a larger size and the percentage of bright phase increased. Based on the EDS result at the coating thickness of 24.9 µm and our previous research about the LWB process of Mg/Cu coated Ti [30], the three phases in the IMC layer were conducted as AlCu2Ti, Ti2Cu and Ti3Al. In addition, microscopic crack was observed in the fusion zone near the IMC/FZ interface as shown in Fig. 7(j), making it a weak zone of IMC/FZ interface. When under tensile load, fracture path was likely to cross IMC layer. Similar fracture path across IMC layer was also referred in the case of Al/steel butt joint using 4047 AlSi12 filler metals [32]. At the lower interface (Fig. 7(k)), a thick and uneven IMC layer was also observed due to the increasing amount of Cu atoms involved in fusion zone. Although the interfacial metallurgical reaction enhanced the bonding strength of the interface to some extent, the microscopic crack near the upper interface was considered as a dominating factor of the bonding strength.

12.42 at% Al and 58.32 at% Ti was identified as AlCu2Ti. The distribution of IMC layer was homogeneous and the IMC/FZ interface was smooth. In addition, there were not any micro-cracks in the IMC layer. The layer II was mainly composed of Ti3Al by EDS results (P3). Due to the interfacial metallurgical reaction, the voids close to interface at the coating thickness less than 10.8 μm disappeared and a tight bonding between layer II and FZ was observed increasing the joint strength to some extent. As for the lower interface as shown in Fig. 7(e), a mechanical bonding without IMC layer was still observed. In the FZ, the grey phase (Mg17(Al,Cu)12) tended to grow up. This was because more Al element in Mg17Al12 was replaced by the increasing addition of Cu element. Fig. 8 shows the EDS mapping results at the upper interface when the coating thickness was 19.7 µm. Al element was enriched uniformly at the interface (Fig. 8(d)) confirming the detection of layer II (Ti3Al) shown in Fig. 7(c). Additionally, the uneven distribution of Cu element (Fig. 8(e)) was observed in layer I, which confirmed the composition of IMC layer at the upper interface (Ti3Al and AlCu2Ti in Fig. 7(d)). Fig. 7(g)–(i) shows the morphologies with SEM back-scattering when the Cu coating thickness was 24.9 μm. At the upper interface (Fig. 7(g)), with more Cu atoms accumulated at Mg/Ti interface, Ti3Al transition layer disappeared and ternary IMC layer bonded to the FZ directly. The shape of IMC/FZ interface was irregular with some sharp corners. IMC layer with excessive thickness was produced in this case, resulting in an embrittlement tendency. The cracks were observed near the interface, which became the origin of the fracture. The microstructure of the interface evolved into a snowflake-like structure composed of three kinds of phase, as shown in Fig. 7(i). The brightest phase (P7) was confirmed as AlCu2Ti, the grey phase (P8) was Ti2Cu and the dark phase (P9) was Ti3Al by the EDS results. To further confirm the phase structure formed in IMC layer at the upper interface, micro-XRD test was employed along the IMC layer and the result was shown in Fig. 9. Besides α-Ti and α-Mg, diffraction peaks of Ti2Cu, Ti3Al and AlCu2Ti were observed, verifying the EDS result mentioned above. As for the lower interface as shown in Fig. 7(h), no IMC layer was still observed. In the FZ close to the lower interface, the original grey phase disappeared. Meanwhile, a new network structure (P6) was observed, which was inferred as eutectic structure (Mg2Cu + α-Mg) based on EDS result. Such a similar result was also reported in our previous research of LWB process of Mg/Cu joint [31]. Fig. 10 shows the EDS mapping results when the coating thickness was 24.9 µm at the upper interface. Much more Cu element was observed in IMC layer with thicker Cu coating, compared with EDS

4.3. Thermodynamic calculation It was worth noticing that Ti-Al intermetallic compound was not observed until the coating thickness was over 19.7 µm under the constant supplementation of Al element, indicating that the addition of Cu had an influence on the metallurgical reaction of Ti and Al. The influence of Cu on formation of Ti3Al intermetallic compound was calculated through thermodynamic calculation. Additionally, the formation Ti2Cu and AlCu2Ti intermetallic compound layer was also analyzed through thermodynamic calculation. The finite element model (FEM) was developed to simulate the temperature field of the joint, providing a calculation temperature T used for thermodynamic calculation. The temperature field of the joint and thermal cycle at upper interface and lower interface of the joint are shown in Fig. 11. Through the result calculated by FEM, the peak temperature of the upper interface was 1265 °C. As for the lower interface the peak temperature decreased to 1202 °C. According to the phase diagrams of Al-Ti, Cu-Ti and Al-Cu-Ti, the formation temperatures of Ti3Al, AlCu2Ti and Ti2Cu were 1180 °C, 965 °C and 960 °C, respectively. It indicated that all the phases mentioned above could be produced at the upper and lower interface, from the viewpoint of 205

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Fig. 7. SEM microstructure of laser welded-brazed Mg/Cu coated Ti joints at different zones with different Cu coating thickness δ: (a)–(c) 10.8 μm, (d)–(f) 19.7 μm, (g)–(i) 24.9 μm; (j)–(l) 28.2 μm.

element were ignored due to their tiny amounts in liquid pool. To analyze the atomic diffusion behavior of Cu, Ti, Mg and Al during the welding process, the dissimilar joint was divided into three region as shown in Fig. 12(a): the upper Mg/Ti interface (region I), the lower Mg/ Ti interface (region II) and the weld seam (region III). At the Mg/Ti interface (region I and II), AZ92-Ti-Cu pseudo-ternary system was established. In this system, the content of Al element was set as 10.1 at%, converting through the weight ratio of 9% in weld filler. As for the weld

formation temperature. A thermodynamic calculation was conducted for a further analysis. For a more convenient calculation of GID, the calculated temperature of 1265 °C and 1202 °C through FEM were taken place of by 1500 K. Formation enthalpy, chemical potential and Gibbs free energy of binary system calculated by Toop model and chemical potential of the elements in Cu-Ti-Mg-Al quaternary system calculated by the extended Toop model were discussed. The influence of Si, Mn, Zn and other 206

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Table 4 Component of different positions in Fig. 7. Position

Mg

Al

Cu

1 2 3 4 5 6 7 8 9

65.36 78.85

34.64 11.43 19.97 10.68 12.42 8.96 14.29 12.63 17.19

9.72 2.57 5.97 26.58 11.97 35.28 29.39 10.63

1.12 1.70 79.07 1.22 2.41 1.41

Ti

77.46 82.23 58.32 48.32 54.54 69.89

Possible phases Mg17Al12 Mg17(Al,Cu)12 Ti3Al Ti3Al AlCu2Ti (α-Mg + Mg2Cu) AlCu2Ti Ti2Cu Ti3Al

seam (region III), since the content of Ti element was closed to zero according to the EDS result of Fig. 7, the Mg-Al-Cu ternary system was established to investigate the reaction mechanism of intermetallic compound in weld seam. Fig. 12(b) shows the formation enthalpies of Mg-Al-Cu ternary system. The Mg atoms in molten weld seam were most likely to react with the Cu atoms in AZ92 weld filler. However, in the case of low additive amount of Cu element, it was difficult for Mg atoms to formed intermetallic compound with Cu atoms in weld seam. As a result, the Mg17Al12 was produced finally when the Cu coating thickness was less than 19.7 μm. With the increase of coating thickness, more and more Cu atoms was involved into weld seam. According to the thermodynamic analysis result mentioned above, the excess Cu atoms would replace some of the Al atoms producing Mg17(Al,Cu)12 during cooling process. The chemical potential of Cu in Mg-Al-Cu system at 1500 K shown in Fig. 12(c) suggested that the Cu atoms tended to diffused to Al atoms, which made it easier for the generation of Mg17(Al,Cu)12. When the Cu coating thickness increased to some extent, the eutectic structures (Mg2Cu + α-Mg) were produced. In addition, Fig. 12(d) shows the Gibbs free energy in Mg-Al-Cu ternary system at the temperature of 1500 K. In the region III, the Cu content was still at a relatively low value (less than 20%) based on the EDS result, as a result it was not necessary to discuss the condition of the Cu content over 20%. Meanwhile, the minimum content of Mg atoms was 65.36% only the region that the content of Mg was over 65.36% marked with dotted borders in

Fig. 9. Micro-XRD pattern along the interface of Mg/Cu coated Ti at the upper interface produced with 24.9 µm coating.

Fig. 12(b) was investigated. In this region, the Gibbs free energy deceased with the increasing Cu content when the Cu content was less than 20%, which could confirm the result that the amount of intermetallic compound produced by Mg, Al and Cu atoms increased obviously with the thicker Cu coating. The calculation result was shown in Table 5 in the Mg content at 70.1% for example. Fig. 13 shows the calculation result of thermodynamic analysis at the upper and lower interface (region I and II). From the formation enthalpy of ternary system of Ti-Al-Cu when the temperature was 1500 K (Fig. 13(a)), the standard molar enthalpy of Mg-Ti was positive, which suggested that no spontaneous reaction occurred between Mg atoms and Ti atoms [24]. While, the lowest formation enthalpy in Ti-AlCu ternary system occurred between Al and Ti atoms, suggesting that Al-Ti phase was much easier to precipitate firstly during the cooling process. Besides, the formation enthalpy of AlCu2Ti was lower than that of Ti2Cu. It suggested that the AlCu2Ti was easier to produce than Ti2Cu, which matched the inference about the IMC layer with various Cu coating thickness as shown in Fig. 7. In AZ92-Ti-Cu system the chemical potential of Cu was shown in Fig. 13(b). To study the

Fig. 8. Elemental distribution of Mg/Cu coated Ti interface at the upper interface with the coating thickness of 19.7 μm: (a) SEM micrograph of interfacial microstructure; (b)–(e) Mg, Ti, Al, Cu EDS mapping result. 207

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Fig. 10. Elemental distribution of Mg/Cu coated Ti interface at the upper interface with the coating thickness of 24.9 μm: (a) SEM micrograph of interfacial microstructure; (b)–(e) Mg, Ti, Al, Cu EDS mapping result.

process of Cu atoms was shown in Fig. 13(d). Only after the diffusion from layer III to layer IV, the Cu atoms could diffuse to the Ti interface directly. Thus, with a Cu molar fraction higher than 0.451 at the upper interface, Cu atoms directly diffused to the Ti substrate, corresponding to the straight black arrows in Fig. 13(b). However, for the lower interface, the original Cu atoms was much fewer than that at the upper interface, which corresponded to the red arrows shown in Fig. 13(b). At a lower molar fraction, taking 0.101 Cu for example, Cu atoms could only diffuse from P1 to P2 corresponding to the process from layer I to layer II shown in Fig. 13(d). As more and more Cu atoms diffused to layer II, the content of Cu gradually increased and reached the content at P3. Then, Cu atoms at layer II (P3) tended to diffuse to layer III (P4) under the impetus of chemical potential. By that analogy, the Cu atom diffused to the Ti substrate intermittently and very slowly, from layer I to layer II to layer III and finally to layer IV, as shown in Fig. 13(d). The enrichment of Cu element at the upper interface made it possible for the generation of Cu-Ti and Al-Cu-Ti (Ti2Cu and AlCu2Ti) at the interface when the coating thickness was greater than 19.7 µm. While for the lower interface, the diffusion of Cu atoms was slow and complex, resulting in a low-Cu-content area at the interface. This Cu content was not enough for the formation of Cu-Ti and Al-Cu-Ti compound, and this was why they were not observed along the lower interface and in the FZ close to Ti as shown in Fig. 7(e). It could also confirm the fact that the amount of Mg-Cu and Mg-Al-Cu compound at the lower interface was much higher than that at the upper interface, as shown in Fig. 7. In addition, when the coating thickness was over 28.2 µm, at the lower interface the IMC layer was also discovered due to the excessive addition of Cu atoms. To explore why Al-Ti intermetallic compound appeared and grew with increase of Cu coating thickness, the chemical potential of Al and Ti were calculated through Ti-Cu-AZ92 system when the temperature was 1500 K, as shown in Fig. 13(e) and (f). As indicated by the black arrows, Al tended to diffuse to Ti and Cu, while Ti preferred to diffuse to Cu, because of the driving force of decreased chemical potential. During diffusion process, Cu atoms played a role in attracting Ti and Al atoms, which would promote the metallurgical reaction of Al and Ti. Such diffusion trend explained that Al element enriched at the interface gradually as the Cu coating became thicker, as shown in EDS mapping

Fig. 11. The thermal cycle curve of typical points at the interface.

influence of different Cu coating thickness on the distribution of Cu atoms, the lowest value of Cu chemical potential element with various molar fraction of Cu was marked with blue crosses and the exact value was shown in Table 6. The characteristic of Cu diffusion was divided into two parts to study based on the turning point at 0.451 Cu. With Cu molar fraction higher than 0.451, Cu atoms would diffuse to Ti substrate directly. While for Cu molar fraction lower than 0.451, Cu atoms diffused to Ti substrate intermittently, as discussed in our previous report [32]. Fig. 13(c) shows the flow field of Mg/Ti butt joint based on a previous research [33], which was helpful to analyze the distribution of the molten Cu coating. An inner circulation was observed at the upper interface. Part of the Cu atoms in molten pool would flow back to the upper interface indicating a constant supply of Cu atoms. Such occasion matched the actual situation with a high Cu molar fraction (maybe higher than 0.451) at upper interface, resulting in obvious interfacial layer. As for the lower interface, the most molten Cu coating was involved into the seam, as shown in Fig. 13(c). It corresponded to the situation at the Cu molar fraction lower than 0.451, which was much lower than that at the upper interface. The microscopic diffusion 208

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Fig. 12. Results of thermodynamic calculation at 1500 K: (a) element distribution schematic of different regions; (b) formation enthalpies of Mg-Al-Cu ternary system; (c) Gibbs free energy of Mg-Al-Cu ternary system; (d) chemical potential of Cu element of Mg-Al-Cu system.

correspondingly without the attraction effect of Cu atoms, as shown in Fig. 14(c). During cooling process, when the temperature was higher than 1180 °C, the formation temperature of metallic compounds in Mg-Al-TiCu quaternary system was not satisfied. Then, when the temperature dropped to 1180 °C, Ti3Al formed firstly at the upper interface. At the temperature of 965 °C with the coating thickness of 19.7 µm, Cu atoms took part in interfacial reaction and AlCu2Ti was produced as shown in Fig. 14(d) [34]. As a result, a compact and smooth IMC layer including Ti3Al and AlCu2Ti was produced. The case of Cu coating thickness over 24.9 µm with the temperature under 960 °C was shown in Fig. 14(f)–(h). Ti2Cu was produced with reaction between the Ti atoms diffusing from the interface and the excessive Cu atoms after the formation of AlCu2Ti [35]. As a result, an IMC layer including Ti3Al, AlCu2Ti and Ti2Cu was produced. In addition, due to the excessive supply of Cu, the thickness of IMC layer increased immoderately and the IMC/ FZ interface became uneven. When the temperature dropped to room temperature, the microstructure evolution in FZ was studied as shown in Fig. 14(g)–(i). In the case of upper interface at the coating thickness of 19.7 µm, most of Cu atoms was involved into weld seam, taking part in the interfacial reaction. A small number of residual Cu atoms replaced some Al atoms in Mg17Al12 and produced Mg17(Al,Cu)12. As a result, Mg17Al12 and Mg17(Al,Cu)12 were observed in FZ. With regard to the coating thickness over 24.9 µm, the amount of the residual Cu atoms increased and reacted with Mg atoms, which produced the (α-Mg + Mg2Cu) eutectic, as shown in Fig. 14(h). As for the lower interface, according to the result of thermodynamic calculation shown in Fig. 13(d), it was difficult for the Cu atoms to get close to the Ti interface and few Cu atoms took part in the interfacial reaction. Because of that, the amount of Cu atoms was high in the weld seam closed to interface. As a result, a large amount of eutectic structure (α-Mg + Mg2Cu) phase was observed near the interface, as shown in Fig. 14(i).

Table 5 The Gibbs free energy of the different Cu content with the Mg content of 0.701 at the temperature of 1500 K (molar fraction). Mg

Al

Cu

Gm (KJ/mol)

0.701 0.701 0.701 0.701 0.701 0.701

0.291 0.251 0.211 0.171 0.131 0.091

0.008 0.048 0.088 0.128 0.168 0.208

−98.20441 −99.09391 −99.76726 −100.283 −100.649 −100.861

result when the coating thickness was 19.7 µm (Fig. 8) and 24.9 µm (Fig. 9). Once the Cu molar fraction was high enough, the Al-Ti metallic compound with the lowest formation enthalpy was produced preferentially during the cooling process, according to the thermodynamic result of binary system. 4.4. Joining mechanism Based on the microstructural evolution characteristics, the phase structure of reaction products and the result of thermodynamic calculation, the joining mechanism was discussed. The schematic of joining mechanism of the welding process was shown in Fig. 14. During the process of heating and atomic diffusion, heat transferred through the molten filler. Cu coating melted and was involved into weld seam, as shown in Fig. 14(a). At the upper interface, according to the calculation result of flow field (Fig. 13(c)), the inner circulation would give a constant supply of Cu atoms. The enrichment of Cu atoms at upper interface played a role in attracting Ti and Al atoms, resulting in the enrichment of both Al and Ti atoms as shown in Fig. 14(b). While at the lower interface, Cu atoms were taken away with the molten filler into the weld seam leading to few Cu atoms concentrated at the lower interface. Meanwhile, the amount of Al and Ti atoms decreased 209

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Fig. 13. Results of thermodynamic calculation at 1500 K: (a) formation enthalpies of Ti-Al-Cu ternary system; (b) chemical potential of Cu in the Ti-Cu-AZ92 pseudoternary system; (c) flow flied of distribution melted Cu coating; (d) schematic of Cu atoms diffusion process on interface; (e) and (f) chemical potential of Ti and Al in the Ti-Cu-AZ92 pseudo-ternary system.

4.5. Mechanical properties

Table 6 The position of the lowest chemical potential of Cu element with different Cu content at the temperature of 1500 K (molar fraction). Cu

Mg

Ti

μCu (KJ/mol)

0.101 0.201 0.301 0.401 0.451 0.501 0.601 0.701 0.801

0.351 0.251 0.151 0.051 0.001 0.001 0.001 0.001 0.001

0.447 0.447 0.447 0.447 0.447 0.397 0.297 0.197 0.097

−130.55 −116.96 −107.92 −101.04 −98.13 −95.340 −89.981 −84.814 −79.780

The fracture load of tensile specimen with width of 10 mm welded through various coating thickness was shown in Fig. 15. The fracture load without Cu coating was lower than that obtained with Cu coating. The low value without Cu coating was mainly attributed to the insufficient spread of molten filler and mechanical bonding. The similar result was also found in our research of Mg/Ti LWB process via Ni coating [24]. When the coating thickness changed to 10.8 μm the fracture load reached 2557 N, which was caused by the improvement of the wettability of molten AZ92 filler on titanium base metal (Figs. 5 and 6). Until the coating thickness reached 19.7 μm, the fracture load still increased, with the dominant factor changing into interfacial reaction. When the coating thickness was 19.7 μm, the joint strength reached the maximum value of 3457 N, which was 85.35% of that of Mg base metal. However, when the coating thickness exceeded 19.7 μm, the fracture 210

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Fig. 14. Schematic of the bonding mechanism: (a) heating process; (b)–(c) diffusion process of Ti atoms, Al atoms and Cu atoms at the upper interface and lower interface; (d)–(i) solidification behavior at the upper and lower interface of the Mg/Cu coated Ti joint with decreasing temperature.

load decreased apparently. In this case, IMC layer became thicker and increased embrittlement inducing formation of micro-crack. According to our previous research about Mg/Ti dissimilar joint [23], the joint fracture morphology correlated intimately to the coating thickness. The SEM micrograph and main fracture mode were shown in Fig. 16, Table 7 shows the corresponding EDS results. The failure the dissimilar joint could be divided into two modes: FZ fracture and interfacial failure, which was shown in Fig. 16(a), (d) and (i). As shown in Fig. 16(a)–(c), when the coating thickness was 10.8 μm, the fracture mode of interfacial failure was observed. Tear ridge was observed at Mg side. As for Ti side the fracture surface nearly completely consisted of large smooth areas, which always presented low joint strength. According to the EDS analysis result at Mg side, the bright particle (#1) was confirmed as Mg17Al12 and the grey phase (#2) was confirmed as α-Mg. As shown in Fig. 16(d)–(h), when the coating thickness was

Fig. 15. Fracture loads of Mg/Cu coated Ti joints versus Cu coating thicknesses. 211

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Fig. 16. Fracture surface of Mg/Cu coated Ti with various coating thicknesses: (a)–(c) δ = 10.8 μm; (d)–(h) δ = 19.7 μm; (i)–(k) δ = 24.9 μm.

19.7 μm, the fracture mode of FZ fracture was observed. In this case the fracture surface was divided into two zones as shown in Fig. 16(e): the interface fracture zone (Fig. 16(f)–(g)) and weld seam fracture zone (Fig. 16(h)). Fig. 16(f) and Fig. 16(g) show the fracture surfaces of interface zone at Ti side and Mg side, respectively. The interface fracture zone changed from tear ridge and smooth areas to widespread scraggly areas, which indicated the enhancement of joint according to the existing literature [16]. In addition, the fracture surfaces on both Ti and Mg sides presented a similar morphology, which was attributed to the fracture of IMC layer. According to the EDS result (#3) on Ti side, the lamellar phase was identified as AlCu2Ti ternary phase, which was

Table 7 Component of different positions in Fig. 16. Position

Mg

Al

1 2 3 4 5

72.36 89.82.

27.64 10.18 15.83 8.18 11.95

9.55 58.73

Cu

18.08 14.76 12.33

Ti

Possible phases

66.09 66.85 17.00

Mg17Al12 α-Mg AlCu2Ti Ti2Cu AlCu2Ti

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closely related to Cu coating thickness. The thin coating thickness (< 19.7 µm) resulted in interface failure, while the thick coating (> 19.7 µm) thickness caused FZ fracture. The maximum fracture load of 3457 N was obtained when the coating thickness was 19.7 μm, representing 85.35% of the Mg base metal. Acknowledgments The study was supported in National Natural Science Foundation of China (Grant No. 51875129) and finance by National Key R&D Program of China (Grant No. 2018YFB1107900). References [1] B. Mansoor, S. Mukherjee, A. Ghosh, Microstructure and porosity in thixomolded Mg alloys and minimizing adverse effects on formability, Mater. Sci. Eng., A 512 (2009) 10–18. [2] M.K. Kulekci, Magnesium and its alloys applications in automotive industry, Int. J. Adv. Manuf. Technol. 39 (2008) 851–865. [3] L. Li, C. Tan, Y. Chen, et al., CO2 laser welding–brazing characteristics of dissimilar metals AZ31B Mg alloy to Zn coated dual phase steel with Mg based filler, J. Mater. Process. Technol. 213 (2013) 361–375. [4] M. Peters, J. Kumpfert, C.H. Ward, et al., Titanium alloys for aerospace applications, Adv. Eng. Mater. 5 (1995) 419–427. [5] W. Moćko, A. Brodecki, Application of optical field analysis of tensile tests for calibration of the Rusinek-Klepaczko constitutive relation of Ti-6Al-4V titanium alloy, Mater. Des. 88 (2015) 320–330. [6] Anas M. Atieh, Tahir I. Khan, Effect of process parameters on semi-solid TLP bonding of Ti–6Al–4V to Mg–AZ3, J. Mater. Sci. 48 (2013) 6737–6745. [7] M.M. Schwartz, Handbook of Metal Welding, National Defense Industrial Press, Beijing, 1988. [8] A.M. Atieh, T.I. Khan, Effect of interlayer thickness on joint formation between Ti6Al-4V and Mg-AZ31 alloys, J. Mater. Eng. Perform. 23 (2014) 4042–4054. [9] A.M. Atieh, T.I. Khan, Application of Ni and Cu nanoparticles in transient liquid phase (TLP) bonding of Ti-6Al-4V and Mg-AZ31 alloys, J. Mater. Sci. 49 (2014) 7648–7658. [10] A.M. Atieh, T.I. Khan, TLP bonding of Ti-6Al-4V and Mg-AZ31 alloys using pure Ni electro-deposited coats, J. Mater. Process. Technol. 214 (2014) 3158–3168. [11] A.M. Atieh, T.I. Khan, Transient liquid phase (TLP) brazing of Mg-AZ31 and Ti-6Al4V using Ni and Cu sandwich foils, Sci. Technol. Weld. Joining 19 (2014) 333–342. [12] M. Aonuma, K. Nakata, Effect of alloying elements on interface microstructure of Mg-Al-Zn magnesium alloys and titanium joint by friction stir welding, Mater. Sci. Eng., B 161 (2009) 46–49. [13] M. Aonuma, K. Nakata, Effect of calcium on intermetallic compound layer at interface of calcium added magnesium-aluminum alloy and titanium joint by friction stir welding, Mater. Sci. Eng., B 173 (2010) 135–138. [14] R. Cao, T. Wang, C. Wang, et al., Cold metal transfer welding–brazing of pure titanium TA2 to magnesium alloy AZ31B, J. Alloy. Compd. 605 (2014) 12–20. [15] C. Xu, G. Sheng, Y. Deng, et al., Microstructure and mechanical properties of tungsten inert gas welded–brazed Mg/Ti lap joints, Sci. Technol. Weld. Joining 19 (2014) 443–450. [16] C. Xu, G. Sheng, H. Wang, et al., Reinforcement of Mg/Ti joints using ultrasonic assisted tungsten inert gas welding brazing technology, Sci. Technol. Weld. Joining 19 (2014) 703–707. [17] M. Gao, Z.M. Wang, X.Y. Zeng, Laser keyhole welding of dissimilar Ti-6Al-4V titanium alloy to AZ31B magnesium alloy, Metall. Mater. Trans. A 43 (2012) 163–172. [18] M. Gao, Z.M. Wang, J. Yan, et al., Dissimilar Ti/Mg alloy butt welding by fibre laser with Mg filler wire-preliminary study, Sci. Technol. Weld. Joining 16 (2013) 488–496. [19] M. Kimura, H. Ishii, M. Kusaka, et al., Joining phenomena and joint strength of friction welded joint between aluminium–magnesium alloy (AA5052) and low carbon steel, Sci. Technol. Weld. Joining 14 (2009) 655–661. [20] S. Lin, J. Song, C. Yang, et al., Metallurgical and mechanical investigations of aluminium-steel butt joint made by tungsten inert gas welding-brazing, Sci. Technol. Weld. Joining 14 (2013) 636–639. [21] S. Chen, L. Li, Y. Chen, et al., Si diffusion behavior during laser welding-brazing of Al alloy and Ti alloy with Al-12Si filler wire, Trans. Nonferrous Met. Soc. China 20 (2010) 64–70. [22] S. Chen, L. Li, Y. Chen, et al., Joining mechanism of Ti/Al dissimilar alloys during laser welding-brazing process, J. Alloy. Compd. 509 (2011) 891–898. [23] C. Tan, J. Yang, X. Zhao, et al., Influence of Ni coating on interfacial reactions and mechanical properties in laser welding-brazing of Mg/Ti butt joint, J. Alloy. Compd. 764 (2018) 186–201. [24] C. Tan, Q. Lu, B. Chen, et al., Influence of laser power on microstructure and mechanical properties of laser welded-brazed Mg to Ni coated Ti alloys, Opt. Laser Technol. 89 (2017) 156–167. [25] C. Zang, J. Liu, C. Tan, et al., Laser conduction welding characteristics of dissimilar metals Mg/Ti with Al interlayer, J. Manuf. Processes 32 (2018) 595–605. [26] K. Zhang, J. Liu, C. Tan, et al., Dissimilar joining of AZ31B Mg alloy to Ni-coated Ti6Al-4V by laser heat-conduction welding process, J. Manuf. Processes 34 (2018)

Fig. 17. XRD result of fracture surface at Mg side with the coating thickness of 24.9 µm.

almost the same as the result of IMC layer at the interface (Fig. 7(d)). It further verified the fact that the fracture path did not propagate along IMC-Ti interface but inside the IMC layer. As for the weld seam fracture zone, on both Mg and Ti side, a typical dimple feature was observed as shown in Fig. 16(h). As the coating thickness increased over 24.9 μm, the fracture mode was still FZ fracture, as shown in Fig. 16(i). At the Ti and Mg side, the lamellar phase composed of Ti2Cu and AlCu2Ti through EDS result on Ti side (#4) and Mg side (#5) became thicker, which was similar with the corresponding result explaining the fracture path cracked inside the IMC layer. Fig. 17 shows XRD result of fracture surface of Mg/Ti at interfacial failure mode at the coating thickness of 24.9 µm on Mg side. The same phase of IMC layer including Ti2Cu, Ti3Al and AlCu2Ti were detected, further confirming that the fracture path propagated in the IMC layer. Besides, it was worth noticing that an obvious crack was discovered at the Ti side as presented in Fig. 16(k), which was ascribed to the embrittlement of IMC layer. In conclusion, the metallurgical reaction was expected to increase the strength of Mg/Ti interface. IMC layer was observed and metallurgical reaction appeared at Mg/Ti interface leading to the increase of joint strength as well with the increase of coating thickness. However, a trend of embrittlement occurred inside IMC layer with increasing Cu content, which made damage to mechanical properties of the dissimilar joint. When the Cu coating was thick to some extent, the fracture path turned from FZ/Ti interface to inside IMC layer. 5. Conclusions (1) The appearances of Mg/Ti butt joints were much improved via the addition of Cu coating, which played a role in improving spreading and wetting influence of molten filler during the stage of metal transfer. (2) The bonding mechanism at the upper interface evolved from a mechanical bonding into a “Ti3Al + AlCu2Ti” metallurgical bonding, and into a “Ti3Al + Ti2Cu + AlCu2Ti” metallurgical bonding with increasing Cu coating thickness. As for lower interface, mechanical bonding evolved into “Ti3Al + Ti2Cu + AlCu2Ti” metallurgical bonding when the coating thickness was over 28.2 µm. (3) It could be confirmed through thermodynamic calculation of Mg-AlCu-Ti system that Cu element played a role in promoting the mutual diffusion of Ti and Al atoms. Ti-Al had the lowest formation enthalpy, which suggested that Ti-Al phase precipitated easily when compared with Al-Cu-Ti and Cu-Ti phase during interfacial metallurgical reaction. (4) The fracture mode could be divided into two types, which were

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