Dissimilar joining of AZ31B Mg alloy to Ni-coated Ti-6Al-4V by laser heat-conduction welding process

Dissimilar joining of AZ31B Mg alloy to Ni-coated Ti-6Al-4V by laser heat-conduction welding process

Journal of Manufacturing Processes 34 (2018) 148–157 Contents lists available at ScienceDirect Journal of Manufacturing Processes journal homepage: ...

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Journal of Manufacturing Processes 34 (2018) 148–157

Contents lists available at ScienceDirect

Journal of Manufacturing Processes journal homepage: www.elsevier.com/locate/manpro

Technical Paper

Dissimilar joining of AZ31B Mg alloy to Ni-coated Ti-6Al-4V by laser heatconduction welding process

T



Kaiping Zhanga,b, Jinge Liua, Caiwang Tana,b, , Gang Wanga,c, Xiaoguo Songa,b, Bo Chenb, Liqun Lia, Jicai Fenga,b a

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 c School of Mechanical and Automotive Engineering, Anhui Polytechnic University, 241000, China b

A R T I C LE I N FO

A B S T R A C T

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

To improve the metallurgical bonding of immiscible and nonreactive Mg/Ti system, Ni electro-coating was employed acting as an intermediate element by reacting with both Mg and Ti in laser heat-conduction welding process. Influence of laser power on welding quality was investigated and interfacial reaction with participation of Ni coating was analyzed. The elemental diffusion mechanism was clarified with the assistance of numerical simulation. Reliable joints were obtained with suitable laser powers ranging from 1200 W to 1800 W. Continuous reaction layer of Ti2Ni and Ti solid solution formed at the interface. The reaction layer thickness was not monotone increasing with increase of laser power. Mg-Al-Ni ternary phase dispersed in fusion zone at Mg side. Diffusion of electroplated Ni was influenced by a bidirectional mechanism to Mg side and Ti side, which resulted in the variation of reaction layer thickness. The maximum tensile-shear load could reach 144 N/mm (about 53.3% of joint efficiency of base metal AZ31B Mg alloy) when laser power was 1500 W. Joint strength and fracture mode were associated with two aspects: interfacial reaction and weld appearance of Mg side. As a result, reliable joining was realized based on appropriate interfacial metallurgical bonding and no weld defects at Mg side.

1. Introduction Magnesium (Mg) and its alloys have recently been attracting growing interest due to their low density, high specific strength, high castability and high utilization ratio for recycling. Mg alloys are widely used in automotive, shipbuilding, aerospace and other industries [1–3]. Titanium (Ti) and its alloys have been another kind of important materials in aerospace industry as a 21st-century metal because of high specific strength, high specific stiffness, preferable oxidation resistance and corrosion resistance [4–6]. With the development of lightweight hybrid components in automotive industry and growing demand for different functions in one technological area, reliable joining between Mg alloys and Ti alloys is required. Successful fabrication of Mg/Ti hybrid structural components would in turn expand the materials application scope. However, the joining between Mg alloys and Ti alloys is restricted by their huge differences in physical and metallurgical properties. The melting point of magnesium and titanium is 649 °C and 1668 °C respectively and the boiling point of magnesium is 1090 °C, hence both of



Mg and Ti alloys are difficult to remain in molten state at the same time. The severe evaporation of magnesium would occur when using the conventional fusion welding techniques. In addition, no mutual diffusion or reaction occurs during solidification process due to the metallurgical immiscibility of Mg/Ti. Some intermediate elements are supposed to be added to realize metallurgical bonding of Mg/Ti by producing solid solution or interfacial compounds at interface. Aiming at these huge challenges in joining Mg to Ti, some related researches were performed. Welding techniques including laser beam welding [7–12], cold metal transfer (CMT) welding [13], friction stir welding (FSW) [14–17] and transient liquid phase (TLP) bonding [18–20] have been adopted to join Mg to Ti in previous studies. Intermediate element from base metals, filler metal or coating layers was utilized to improve metallurgical bonding of Mg/Ti immiscible system. As an innovation of arc welding, CMT welding was introduced to join dissimilar metals due to its effective control of heat input [21]. Cao et al. used AZ61-Mg wire (6 wt.% Al) to join pure titanium and AZ31 Mg (3 wt.% Al) through CMT welding. From this study, the Ti3Al, Mg17Al12 phases formed at the brazing interface indicating that element

Corresponding author at: State Key Laboratory of Advanced Welding and Joining, Harbin Institute of Technology, Harbin 150001, China. E-mail addresses: [email protected], [email protected] (C. Tan).

https://doi.org/10.1016/j.jmapro.2018.05.037 Received 27 March 2018; Received in revised form 6 May 2018; Accepted 31 May 2018 1526-6125/ © 2018 Published by Elsevier Ltd on behalf of The Society of Manufacturing Engineers.

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Al from Mg based filler were beneficial to join Mg and Ti [13]. With regard to the feasibility of bonding Mg and Ti using FSW process, Aonuma et al. analyzed the interfacial microstructure of Mg (AZ31)/Ti (pure titanium) FSWed butt joint. The Ti-Al thin layer formed at joint interface, suggesting Al element from Mg base metal played a crucial part in joining Mg to Ti [16]. To enhance the flexibility of adding intermediate element, electrocoating process was employed. Atieh et al. investigated TLP bonding of Mg-AZ31 and Ti-6Al-4 V alloys using pure thin Ni electroplated coating. Two mechanisms were summarized: interdiffusion at the Mg/Ni interface and liquid eutectic formation at interface, as well as solid-state interdiffusion of Ni and Ti [18]. To further study the influence of intermediate elements on interfacial reaction of Mg/Ti joint, they used an electroplated Ni coating containing dispersion of Ni and Cu nanoparticles. Different nanoparticles were found to result in different compounds at the interface. The joint strength was associated with the varied compounds [19]. Laser beam welding had obvious advantages in joining Mg to Ti due to its precise control on heat input. The reliable joining of Mg-AZ31B and Ti-6Al-4 V alloys was realized using laser keyhole welding, and the study found that mechanical properties of Mg/Ti joint was associated with laser heat input and the offset from the laser beam to the edge of weld seam [7]. In our previous studies, the reliable Mg/Ti joint was obtained by laser welding-brazing process. From our researches, element Al played a crucial role adjusting interfacial reaction and the thin reaction layer Ti3Al formed at interface [8]. The electroplated Ni on Ti surface could enhance the wetting-spreading ability of molten filler, and it participated in the interfacial reaction at the same time [10]. The formation of a metallurgical Mg/Ti joint needed some other elements such as Al, Ni or Cu to improve the interfacial reaction greatly. It is also worth noting that dissimilar metals were joined by laser welding in a heat-conduction mode in many other studies such as Al/ steel and Mg/steel [22,23]. Therefore, in this work, laser heat-conduction welding was used to join Mg-AZ31B and Ti-6Al-4 V dissimilar alloys. Electroplated Ni on Ti sheet surface was used as intermediate element to improve the interfacial reaction. The reliable and defect-free Mg/Ti joints with good surface appearances were obtained by adjusting the welding parameters. Subsequently, the process characteristics, interfacial microstructure and mechanical properties were investigated. Based on the above, the diffusion mechanism at the Mg/Ti interface was clarified with the assistance of numerical simulation.

in Fig.1b by scanning electron microscopy (SEM) analysis. Uniform plating process was maintained by magnetic stirring. Before the experiment, grinding was used to remove the oxide from the surface of magnesium alloy. Then the grease and other contaminants on the surface were removed by swabbing with acetone.

2. Experimental

To analyze the welding temperature field and verify the diffusion mechanism, a FE (finite element) model was used in the numerical simulation performed by software MSC. Marc. Large defocused distance was used in the laser heat-conduction welding process in this work, so a gauss plane heat source was adopted as follows:

2.2. Experimental procedure A fiber laser (IPG YLR-6000) possessing a maximum power of 6 kW was employed in this study. The diameter of focused laser beam was 0.2 mm, focused by a 200-mm focal length lens. Fig.2 shows the schematic diagram of laser heat-conduction welding process. A lap configuration by placing Ti alloy sheet on top of Mg alloy sheet was fixed in the assembly of weldments. The laser beam was irradiated on Ni coated surface vertically. Argon shielding gas blowing over the top surface was used to protect the molten pool from oxidation with a flow rate of 20 L/ min. In previous reports [10,24], laser power was found to have conspicuous influence on welding quality of dissimilar metals. In this work, varied laser powers were employed, while the welding speed (1 m/min) and laser defocusing distance (focus distance of positive 30 mm defocused from Ti alloy front surface) were kept constant. For comparison, the joining of Mg to uncoated Ti was also performed although it was difficult to realize the continuous and reliable weld without the addition of element Ni. In this research, the joint of Mg to uncoated Ti with the best properties (obtained at laser power of 1800 W) in our experimental parameter ranges was selected to analyze. Metallographic specimens were prepared by mounting in denture base resins and then grinded by 320, 500, 1000, 1500, 2000 SiC grades abrasive paper and polished a mirror-like surface by 2.5-μm diamond suspension. The weld appearance and microstructures were examined by optical microscopy (OM) and SEM. The distribution of chemical elements across the interface of joint was analyzed through energy dispersive X-ray spectrometer (EDS) analysis. Micro-Vickers’s hardness profile of joint was obtained by a micro-hardness tester under a test load of 100 g and a dwell time of 10 s. The Mg/Ti dissimilar lap joint obtained by laser heat-conduction welding process were cut into 10 mm wide perpendicularly by electro-discharge machining as the tensile test specimens. Then the tensile-shear test was carried out with a cross-head speed of 1 mm/min at room temperature. 3. Numerical simulation

2.1. Materials Commercially available AZ31B Mg alloy and Ti-6Al-4 V Ti alloy with the dimension 100 × 50 × 1 mm were used in the experiment. The supply state of Mg alloy was H24 with annealing. The chemical compositions of AZ31B Mg alloy sheet and Ti-6Al-4 V alloy sheet are listed in the Table 1. The Ti-6Al-4 V alloys were cleaned in acid (15% HCl, 5% HF, 80% distilled water) for 3 min and then rinsed by water to remove surface oxides. Fig.1a shows the electroplating process. The plating solution used for electroplating Ni coating was prepared by dissolving 180 g NiSO4·6H2O, 10 g NaCl, 70 g Na2SO4, 30 g H3BO3 and 60 g MgSO4·7H2O in 1 L of distilled water. The electroplating temperature was kept at 35 °C and the electroplating time was set as 2 h. The thickness of Ni layer on Ti surface was approximately 20 μm shown

qs (x , y ) =

Where, qs (x,y) is linear energy. α is heat efficiency of welding process. Qs is heat source power. rs is the effective radius of the heat source. The FE mesh was set up according to experimental situation shown in Fig. 3. A non-uniform mesh was used in this work, in which a finer mesh with about 0.166 × 0.166 × 0.15 mm3 was used at the molten pool zone and HAZ, so that the numerical precision was guaranteed. A coarse mesh was located at the zone which was far from heat affected zone (HAZ) for numerical simplification. The total numbers of nodes and meshes were 358,490 and 313200, respectively. In this simulation, the environment temperature was set as 20 °C. The numerical simulation model was verified by contrasting the experimental and numerical results shown in Fig. 4. The width of fusion zone at Mg alloy side obtained at laser power of 1500 W was 9.65 mm in experiment which could be observed after etching specimens in a solution of nitric acid and alcohol (95% ethanol and 5% HNO3) for 15 s. The width of fusion

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

AZ31B Ti-6Al-4V

Al

Zn

Mn

Fe

V

Si

Mg

Ti

2.5–3.5 5.5–6.8

0.5–1.5 3.5–4.5

0.2-0.5 –

< 0.005 0.3

– 3.5–4.5

0.1 –

Bal. –

– Bal.

α (x 2 + y 2 ) ⎤ αQs exp ⎡− 2 ⎥ ⎢ πrs rs2 ⎦ ⎣

149

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Fig. 1. The schematic diagram of electroplating and morphology of Ni coating layer: (a) electroplating process, (b) morphology of Ni coating.

Fig. 2. Schematic diagram of laser heat-conduction welding of Mg to Ni coated Ti.

Fig. 3. Finite element analysis model of laser heat-conduction welded Mg/Ti joint.

Fig. 5. Front surfaces of Mg/Ti joints obtained at different laser powers: (a–f) laser power 800 W–2000 W, with Ni coating; (g) without Ni coating.

4. Results and discussion 4.1. Weld appearance Fig. 5 shows typical front surfaces of laser heat-conduction welded Mg/Ti dissimilar joints with or without Ni coating obtained at different laser powers (800 W–2000 W). Instable appearance is presented in Fig. 5a with the laser power of 800 W. The acceptable weld surfaces without distinct defects were obtained using higher laser powers. Front surfaces became smooth and uniform with the increase of heat input

Fig. 4. Comparison of fusion zone at Mg alloy side between experimental and numerical results.

zone at Mg side in numerical results under the identical parameters was 9.63 mm, which was in good agreement with experimental results. The error range of simulation could be controlled within 5% in this work. 150

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used for analyzing the behavior of liquid Mg, which was also associated with joint strength. Completely no bonding joint section was presented in Fig. 6a with laser power of 800 W. Fig. 6b shows a part of bonding between the base metals. The effective joining length enhanced when laser power reached 1200 W presented in Fig. 6c, but the lack of bonding also appeared at this joint. Low heat input was considered as the reason for distinct lack of bonding with laser powers of 1000 W and 1200 W. The sound joint section without obvious defects was shown in Fig. 6d obtained at laser power of 1500 W, and the effective joining length was longest at this welding parameter. However, lack of bonding and lack of Mg base metal appeared simultaneously at joint section as seen in Fig. 6e when laser power was 1800 W. With increase of laser power, the liquid Mg flowed into middle position of cross section more sharply due to action of gravity, and the evaporation of Mg was drastic at the same time. When laser power was 2000 W, lack of bonding and lack of Mg base metal resulted in over-burning (Fig. 6f) even the welding through. The weld defects and lack of effective joining length were harmful to joint strength. As seen in Fig. 6g, without the addition of Ni element, the effective joining length became more insufficient due to the immiscible and nonreactive Mg/Ti. 4.2. Fusion zone and interfacial reaction layer Fig. 7a–e shows the SEM morphologies of interfacial microstructures in Mg/Ni coated Ti joints obtained at different laser powers. Dot-like structures distributed dispersedly in fusion zone of Mg alloy side. This dot-like structure was identified as Mg-Al-Ni ternary phase since it was composed of 55.03 at.% Mg, 23.13 at.% Al, 0.30 at.% Ti and 21.30 at.% Ni according to the EDS analysis result in Table 2 (P1 and P2 marked in Fig. 7). Similar phase was confirmed in our previous work about Mg/Ti welding [10]. It implied that element Ni from electroplated coating diffused and dissolved into molten Mg alloy, and then segregated with element Al producing Mg-Al-Ni ternary phase during solidification. As a result, the diffusion of Ni into Mg fusion zone could reduce reaction layer thickness at Ti alloy side. At the same time, element Ni from pure Ni coating diffused into Ti alloy side and reacted with Ti producing reaction layer. Light structures and dark structures were observed in reaction layer. Light structure was identified as Ti2Ni since it was composed of 1.99 at.% Mg, 4.50 at.% Al, 67.50 at.% Ti and 26.01 at.% Ni by EDS analysis shown in Table 2 (P3 and P4 marked in Fig. 7). The similar observation was reported in previous study about Ti/Ni welding [25]. Based on Ti-Ni phase diagram seen in Fig. 8 [26], element Ni diffused into Ti side, resulting in the formation of Ti solid solution and Ti2Ni. Two ways were supposed to

Fig. 6. Cross sections of Mg/Ti joints obtained at different laser powers: (a–f) laser power 800 W–2000 W, with Ni coating; (g) without Ni coating.

(Fig. 5b–d). Stable and sound welded joints without burning loss are shown in Fig. 5c and d when the laser powers were 1200 W and 1500 W, indicating that heat inputs were appropriate at these parameters. When laser power increased, Ni coatings broke and dissolved into Ti alloys with higher heat input. The pits appeared in front surfaces (Fig. 5e and f), suggesting that welding process became unstable again. As presented in Fig. 6, cross sections of Mg/Ti dissimilar joints were

Fig. 7. Interfacial morphologies of Mg/Ti joints obtained at different laser powers: (a–e) laser power 1000 W–2000 W, with Ni coating; (f) without Ni coating. 151

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Table 2 EDS analysis results of various points marked in Fig. 7 (at.%). Point

Mg

Al

Ni

Ti

Possible phases

P1 P2 P3 P4 P5 P6

55.03 55.48 1.99 0.54 1.53 1.04

23.13 26.13 4.50 5.90 10.50 12.90

21.30 17.83 26.01 30.55 9.74 11.46

0.30 0.42 67.50 63.01 78.23 74.60

Mg-Al-Ni ternary phase Mg-Al-Ni ternary phase Ti2Ni Ti2Ni Ti solid solution + Ti2Ni Ti solid solution + Ti2Ni

Fig. 10. Average thickness of reaction layer in Mg/Ni coated Ti joint interface obtained at different laser powers.

phase diagram to Ti-Al phase diagram (Fig. 9b [29]), the solubility of Al in Ti substrate was higher than that of Ni in Ti substrate because of similarity in atomic size and electronic properties between Al and Ti. Therefore, the element Ni would separate and then transformed into Ti2Ni during cooling process. All analysis indicated that the dark structure was Ti solid solution (containing Al) with precipitated phase Ti2Ni. In a word, the interdiffusion of Ni/Ti caused the formation of these kinds of structures in reaction layer. However, the specific morphologies of the microstructures varied with different laser powers. Note that the thickness of reaction layers was not monotone increasing with increase of laser power as seen in Fig. 10. The reaction layer thicknesses were 21 μm, 38 μm, 8 μm, 19 μm, 34 μm respectively in average with laser power increasing from 1000 W to 2000 W, and the thinnest layer (< 10 μm) formed at laser power of 1500 W (Fig. 7c). Nevertheless, no obvious reaction layer formed at Mg/Ti interface without the addition of Ni element as seen in Fig. 7f. EDS line scanning across interface was performed according to the direction shown in Fig. 7a, and the results are presented in Fig. 11. Obvious reaction layer zone can be observed in Mg/Ni coated Ti joints with all the laser powers. From Fig.11a–e, Mg content decreased sharply and almost disappeared in reaction layer zone, indicating the reaction layer was at Ti side and not associated with diffusion of element Mg. Ti content in reaction layer zone was lower than that in base metal, and then increased gradually from reaction layer to base metal. In reaction layer zone, significant segregation of Ni could be observed,

Fig. 8. Ti-Ni binary phase diagram [26].

cause the formation of Ti2Ni: on one hand, the TiNi phase would first nucleate from the liquid. Next, Ti2Ni tended to precipitate around the TiNi via a peritectic reaction. On the other hand, some other Ti2Ni phases formed during further cooling process because some β-Ti transformed to α-Ti which had lower solid solubility at the lower temperatures. Thus, element Ni would extract from α-Ti resulting in formation of Ti2Ni. The EDS analysis result of dark structure was 1.53 at.% Mg, 10.50 at.% Al, 78.23 at.% Ti and 9.74 at.% Ni shown in Table 2 (P5 and P6 marked in Fig. 7). A relative low-temperature (800 °C) Ti-Ni-Al isothermal section was used for identifying the phase formation of this dark structure as shown in Fig. 9a [27]. Referring to the isothermal section, the composition point of dark structure was located at the region marked with a red line circle in Fig. 9a. Besides, according to the related investigation [28] and the comparison of Ti-Ni

Fig. 9. (a) Ti-Ni-Al isothermal section at 800 °C and (b) Ti-Al binary phase diagram [27,29]. 152

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Fig. 11. Results of line scanning across interface obtained at different laser powers: (a–e) laser power 1000 W–2000 W, with Ni coating; (f) without Ni coating.

Fig. 12. Results of mapping scanning across Mg/Ni coated Ti joint interface: (a) SEM morphology; (b–e) distribution of Mg, Al, Ti and Ni.

thickness in region Ⅱ was much thicker than that in region Ⅰ. This phenomenon would be discussed later.

but the evident enrichments of Al content were only found at laser power of 1500 W and 1800 W. However, no obvious elemental enrichment appeared at Mg/Ti interface without the addition of Ni coating (Fig. 11f). The mapping scanning about reaction layer was used to observe elemental distribution presented in Fig. 12, and the result also showed that some kinds of structures formed in reaction layer. The light structures mainly contained element Ti, Ni and the dark structures mainly contained element Ti, Al and Ni. Fig. 13 presents that the Mg/Ti joint interface was divided into two regions. Region Ⅰ was located in effective joining zone, and the results above mainly focused on this region. Region Ⅱ was near effective joining zone located at the edge of laser direct irradiation zone. Fig. 13b and c show the SEM interfacial morphologies of these two regions, respectively. Separation between Mg and Ti base metals could be seen in Fig. 13c because of low heat input and the gravity of liquid Mg in this region. The microstructures at reaction layer were Ti2Ni and Ti solid solution in region Ⅱ, the same as that in region Ⅰ. The mapping scanning results of region Ⅱ are presented in Fig. 11. Note that the reaction layer

4.3. Numerical simulation results The temperature field on peak temperature and thermal cycle across the Mg/Ti joint were simulated to help clarify the diffusion mechanism. Fig. 14 presents the temperature fields extracted from welding zone marked in a rectangular region in Fig. 6d with three different laser powers (1200 W, 1500 W and 1800 W). The fusion lines were marked representing the melting points of Ni (1455 °C) and Ti (1668 °C) respectively. The states of Ni and Ti at the interface varied with different laser powers. Both Ni and Ti could not reach the melting point at the interface when laser power was 1200 W shown in Fig. 14a and b. The melting of Ni at the interface was observed with laser power of 1500 W shown in Fig. 14c, but the Ti alloy hardly melted at the interface at this case (observed in Fig. 14d). When laser power increased to 1800 W, the obvious melting of Ti alloy occurred at the interface from the 153

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Fig. 13. (a) Optical micrograph of the Mg/Ni coated Ti joint; (b) SEM interfacial morphology of region Ⅰ marked at (a); (c) SEM interfacial morphology of region Ⅱ marked at (a).

simulation result in Fig. 14f, indicating the diffusion and dissolution of Ni/Ti improved markedly. In addition, the thermal cycles at the interface from welding process was extracted (the point marked in Fig. 14a, c, e), and the results are presented in Fig. 15. The peak temperature could not reach the melting point of Ni with laser power of 1200 W. However, the peak temperatures were higher than the melting point of Ni when laser power was 1500 W and 1800 W, revealing that the interdiffusion of Ni/Mg at the interface transformed from liquid/solid to liquid state when laser power increased from 1200 W to 1500 W. The peak temperature was lower than the melting point of Ti at laser power of 1500 W, which was opposite in the case of laser power of 1800 W. It was evident that the interdiffusion mode of Ni/Ti at the interface transformed from liquid/ solid to liquid state when laser power increased from 1500 W to 1800 W.

4.4. Diffusion mechanism of electroplated Ni Fig. 15. Thermal cycles at the interface obtained at different laser powers.

According to the results of numerical simulation, the bidirectional Ni diffusion mechanism during laser heat-conduction welding of Mg to Ti was clarified, and the schematic is shown in Fig. 16. On one hand, element Ni could diffuse into fusion zone at Mg alloy side, and then element Ni segregated with element Al (mainly from AZ31B Mg alloy)

producing Mg-Al-Ni ternary phase. The growth of reaction layer at Ti side was restricted due to element Ni segregation. During welding process, the heat transfer direction was ‘Ti alloy → Ni interlayer → Mg

Fig. 14. Temperature fields at peak temperatures obtained at different laser powers: (a)(b) 1200 W; (c)(d) 1500 W; (e)(f) 1800 W. 154

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Fig. 16. Schematic diagram about Ni diffusion and formation of reaction layer: (a) initial states of base metals; (b) diffusion tendency of element Mg, Al, Ti, Ni; (c–e) diffusion result and phase formation with different laser powers.

and Ti base metals appeared in this region due to low heat input and the gravity of liquid Mg. The Ni diffusion into Mg side became difficult and the heat was blocked at Ti side. More element Ni diffused into Ti side to take part in the formation of reaction layer. As a result, the reaction layer in region IIwas thicker than that in region I at the same laser power.

alloy’ with a lap configuration. In addition, the contact between Ti and Ni was tighter because the interlayer was electroplated directly on the surface of Ti alloy. Therefore, on the other hand, element Ni diffused into Ti alloy side and continuous reaction layer consisting of Ti2Ni and Ti solid solution formed. This process could be observed through Fig. 16b–e. As a result, the more Ni diffusing into Ti alloy side, the thicker reaction layer grew, but the diffusion of Ni into Mg fusion zone could reduce reaction layer thickness. Based on the laser rapid heating and cooling characteristics, the reaction layer thickness was associated with the combined effect of heat input and Ni diffusion. When laser power was 1000 W, the Ni diffusion into Mg side or Ti side was both little, and the average reaction layer thickness was 21 μm shown in Fig. 7a. As laser power rose to 1200 W, the interdiffusion form of Ni/Mg maintained at a solid/liquid state. But the Ni diffusion into Ti side enhanced obviously due to the increase of direct heat input to Ti sheet. As a result, the reaction layer thickness grew to 38 μm shown in Fig. 7b. These two cases above are presented in Fig. 16c. With the laser power of 1500 W, higher heat input could cause melting of Ni coating and more flow of liquid Mg alloy. In this case, the mode of Ni/Mg interdiffusion became more sufficient and a large amount of element Ni mixed into fusion zone rather than participating in the growth of reaction layer. Meanwhile, the heat input could not change the state of Ti alloy resulting in little improvement in Ni/Ti interdiffusion. All these led to the thinnest reaction layer (average < 10 μm) shown in Figs. 7c and 16d. When laser power increased to 1800 W or 2000 W, the interdiffusion mode of Ni/Ti was transformed into liquid-state interdiffusion. Melting of Ti alloy needed greater laser power than 1500 W according to result of numerical simulation, in this case the diffusion of Ni into Ti alloy side became sufficient dramatically with the state change of Ti alloy. The thicker reaction layers were observed in Figs. 7d–e and 16e, since a large amount of element Ni diffused into Ti alloy side. All above was consistent with the simulation results presented in Section of 4.3. The situation in region II (shown in Fig. 13c) was also used for proving the diffusion mechanism above. The separation between Mg

4.5. Mechanical properties The Vickers’s hardness distribution across the interface of Mg/Ni coated Ti joints with varied laser power is shown in Fig. 17. The hardness profile across the interface could reflect the microstructural evolution from Mg alloy fusion zone to Ti alloy side. As can be seen from the figure, the hardness values of Mg and Ti base metals were

Fig. 17. Hardness distribution across the Mg/Ni coated Ti interface obtained at different laser powers. 155

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Fig. 18. Tensile-shear fracture loads with variation of laser power.

Fig. 19. Typical fracture modes of Mg/Ni coated Ti joints with different laser powers: (a) 1000 W; (b) 1500 W; (c) 1800 W.

56 H V and 350 H V in average, respectively. The hardness values in reaction layer zone increased dramatically reaching 595 H V. The hardness indentations at reaction layer zone and Ti base metal side could be seen in the inset of Fig. 17. With the same test load, the areas of indentations at reaction layer were obviously smaller than those at Ti base metal side. The element Ni and Al participated in the formation of solution strengthening and intermetallic compounds, resulting in the higher hardness at the interface. The hardness value obtained in reaction layer zone with laser power of 1000 W was conspicuously lower than others, corresponding with the morphology of microstructures at reaction layer. From Fig. 7a, the hard and brittle light structures were sparser than those obtained with other laser powers. Fig. 18 presents tensile-shear fracture loads of laser heat-conduction welded Mg/Ti joints with or without Ni coating obtained at different laser powers. The fracture load was as low as 107 N/mm without addition of Ni coating due to the weak interfacial bonding and insufficient effective joining length. The fracture load of Mg/Ni coated Ti joint was analyzed as follow. The fracture load was below 115 N/mm when the laser power was 1000 W because the effective joining length was deficient. Fracture load improved sharply with increase of laser power. In the case of 1500 W the fracture load reached a maximum value of 144 N/mm, which was 53.3% of that of Mg base metal. Appropriate interfacial reaction and no welding defects resulted in highest strength. The fracture load then declined with further increase of laser power, since higher heat input would cause welding defects again. The typical fracture modes are exhibited in Fig. 19 and used for explaining the variation of joint fracture load. Crack fractured along the interface with low fracture load at laser power of 1000 W. In this case the heat input was inadequate, which resulted in the insufficient effective joining length and weak metallurgical bonding. These factors caused the lower strength although no obvious defects occurred at Mg side at this parameter. However, the crack did not propagate along the interface in the whole fracture process and the tearing of Mg alloy could be observed in the effective joining region shown in enlarged view of Fig. 19a. It suggested that the slight interfacial reaction occurred in this region. The fracture mode of joints obtained at laser power of 1200 W or the case of no Ni coating was similar to that obtained at laser power of 1000 W, which was not presented here. Fig. 19b presents the fracture mode with laser power of 1500 W. There were also no weld defects at the Mg side, but crack fractured along Mg alloy side rather than along the interface and then propagated far from interface, which was different from the situation in Fig. 19a. Appropriate interfacial reaction resulted in ultra-thin hard reaction layer (< 10 μm) at the interface, which could enhance the strength of this zone. Meanwhile effective joining length was the longest without any defects in this case.

Therefore, the fracture load was the highest at laser power of 1500 W. The third kind of fracture mode obtained at laser power of 1800 W is shown in Fig. 19c, which represented the joints with weld defect about lack of Mg base metal at laser powers of 1800 W and 2000 W. The crack fractured along weld defect in Mg alloy side and then propagated far from the interface in this mode. The increase of heat input and favourable metallurgical bonding made the interface possess enough strength. Nevertheless, the excessive heat input would cause lack of Mg base metal and even welding through, which made the fracture loads of joints obtained at these parameters lower than that obtained at laser power of 1500 W. In summary, the joint strength and fracture mode were attributed to two aspects: interfacial reaction and weld appearance at Mg side. 5. Conclusions 1 Laser heat-conduction welding of Mg to Ni-coated Ti in lap configuration has been successfully achieved. The process parameters were optimized and reliable welding joints were obtained with laser power ranging from 1200 W to 1800 W. 2 Pure Ni coating diffused into Ti alloy side producing reaction layer of Ti solid solution and Ti2Ni during welding process. But the reaction layer thickness was not monotone increasing with the increase of laser power. Mg-Al-Ni ternary phases were distributed in fusion zone of Mg alloy side because of Ni/Mg interdiffusion. 3 A bidirectional Ni diffusion mechanism was clarified based on numerical simulation: With lower heat input, the reaction layer was thicker because the Ni coating remained solid-state during welding process. When the heat input could melt both Mg and Ni but could not melt Ti alloy, element Ni mainly diffused into Mg fusion zone resulting in thinner reaction layer formed at Ti side. In the case of higher heat input, the Ni/Ti interdiffusion transformed from liquid/ solid into liquid state due to melting of both Ni and Ti. A large amount of element Ni diffused into Ti alloy side, thus the reaction layer became thicker again. 4 The maximum tensile-shear force could reach 144 N/mm when laser power was 1500 W, representing 53.3% joint efficiency with respect to Mg base metal. Joint strength and fracture mode were attributed to two aspects: interfacial reaction and weld appearance of Mg side. With higher heat input, joints did not fracture along the interface reflecting that a good metallurgical bonding formed at the interface.

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Acknowledgments The study was financially supported by National Natural Science Foundation of China (Grant Nos. 51504074), Key Research and Development Program of Shandong Province (Grant No. 2017GGX30147 and 2017CXGC0811), China Postdoctoral Science Foundation (No. 2016T90280), and State Key Lab of Advanced Welding and Joining, Harbin Institute of Technology (AWJ-16-M04).

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