316 L stainless steel joint brazed with Ti33.3Zr16.7Cu50−xNix amorphous filler metals

Journal Pre-proof Interfacial microstructure and shear strength of Ti6Al4V alloy/ 316 L stainless steel joint brazed with Ti33.3Zr16.7Cu50−xNix amorphous filler metals

Yueqing Xia, Honggang Dong, Runze Zhang, Yaqiang Wang, Xiaohu Hao, Peng Li, Chuang Dong PII:

S0264-1275(19)30818-4

DOI:

https://doi.org/10.1016/j.matdes.2019.108380

Reference:

JMADE 108380

To appear in:

Materials & Design

Received date:

24 August 2019

Revised date:

19 November 2019

Accepted date:

20 November 2019

Please cite this article as: Y. Xia, H. Dong, R. Zhang, et al., Interfacial microstructure and shear strength of Ti6Al4V alloy/316 L stainless steel joint brazed with Ti33.3Zr16.7Cu50−xNix amorphous filler metals, Materials & Design(2019), https://doi.org/10.1016/j.matdes.2019.108380

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© 2019 Published by Elsevier.

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Interfacial microstructure and shear strength of Ti6Al4V alloy/316L stainless steel joint brazed with Ti33.3Zr16.7Cu50-xNix amorphous filler metals Yueqing Xia, Honggang Dong*, Runze Zhang, Yaqiang Wang, Xiaohu Hao, Peng Li, Chuang Dong School of Materials Science and Engineering, Dalian University of Technology, Dalian 116024, PR China

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Corresponding author. Prof., Ph.D.; Tel.: +86 0411 84706283; E-mail address:

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Abstract

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[email protected] (Honggang Dong)

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Ti33.3Zr16.7Cu50-xNix (x=0-16.5, at.%) amorphous filler metals were designed to

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braze Ti6Al4V alloy and 316L stainless steel (SS). The effect of Ni addition in filler

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metals on the wettability, joint microstructure evolution and shear strength were investigated. The reaction phase identification and crack initiation mechanism were

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analyzed in depth. Ni addition weakened the wettability of the filler metal and thickened the brazed seam. The filler metal with 11 at.% Ni was optimized for brazing of Ti6Al4V alloy/316L SS with the maximum joint shear strength of 318 MPa. FeTi, Fe2Ti, FeCr, and α-Fe phases formed around the transition zone close to 316L SS substrate, which was the weak part of the brazed joints. The interface between FeTi and Fe2Ti phases was non-coherent with the lattice mismatch of 61.4%, initiating the crack at (β-Ti + FeTi)/Fe2Ti interface. The initiative cracks mainly propagated along the Fe2Ti and FeCr layers with brittle feature. Optimizing the constituents and alloying process for Ti-Cu-based filler metal has huge potential in improving the 1

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performance of titanium alloy/steel joint. Keywords: Vacuum brazing; Wettability; Amorphous filler metal; Interfacial microstructure; Shear strength 1 Introduction With excellent corrosion resistance, low cost and weight reduction, hybrid titanium alloy/stainless steel components attracted growing interest in particular

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applications, such as oil pipelines, aerospace and detectors [1, 2]. Both solid-state

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welding [3-5] and fusion welding [6-8] processes have been successfully applied to

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join titanium alloy and stainless steel. However, brazing has more potential for

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welding dissimilar materials. Compared with solid-state welding and fusion welding,

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brazing is more efficient, joint design of titanium alloy/stainless steel is more flexible and the residual stress in the joint is lower.

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Most of the researches paid attention to Ag-based and Ti-based filler metals for

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brazing titanium alloy/stainless steel. Soltani Tashi et al. [9] brazed Ti6Al4V alloy to 316L SS with Ag-Cu-Zn filler metal. The joint shear strength ranged from 45 MPa to 85MPa. Improving the brazing temperature and bonding time increased the amount of detrimental Cu-Ti and Fe-Cu-Ti intermetallic compounds (IMCs). The brazed joint revealed ductile fracture mode at low brazing temperature and time, but brittle feature at high brazing temperature and time. Lee at al. [10] applied Ag-Cu filler and Ag interlayer to join pure titanium/UNS S31254 steel at 820 oC/10 min. The brazed joint presented an interfacial microstructure of titanium substrate/TiAg layer/Ag-rich solid solution/steel substrate. The maximum tensile strength reached 440 MPa with the fracture location at Ag-rich solid solution layer with ductile feature. 2

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Compared with the Ag-based filler metal, the joint brazed with Ti-based filler metal has superior performance like corrosion resistance [11] and high-temperature strength [12]. Bajgholi et al. [13] conducted brazing of Ti6Al4V alloy/316L stainless steel with Zr-Ti-Ni-Cu amorphous filler metal. Only Ni-Ti IMCs formed in the brazed seam, leading to higher micro-hardness than that in substrates. Lee et al. [14] used

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Zr-Ti-Ni-Cu-Be amorphous filler metals to braze titanium (Gr. 2)/UNS S31254 stainless steel. The highest tensile strength was 190 MPa at a lower brazing

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temperature of 800 oC with the formation of β-Ti and NiTi2 phase in the brazed joint.

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Continuous layered brittle NiTi2 and σ phase came into being at the interface of

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brazed seam/steel substrate at 900 °C, which deteriorated the brazed joints. In our

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previous studies [15-17], various amorphous Ti-based filler metals were applied to

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braze Ti6Al4V alloy to 316L stainless steel with lapped joint design. However, low joint strength always occurred due to joint distortion during shear test, which was

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difficult to evaluate the true shear strength of brazed joint. The brazed seam consisted of diffusion zone (β-Ti)/residual filler metal/Ti-Cu-Fe layer/diffusion zone (Fe2Ti and FeCr). All brazed joints fractured along the brazed seam/316L SS substrate interface, where brittle Fe-Ti IMCs formed. While brazing titanium alloy/stainless steel with Ti-based filler metal, researchers currently often focused on the brazing parameters with a single brazing filler metal. However, the effect of filler constituent on the joint performance lacks discussion in depth, and the identification of reaction phases is ambiguous and only relying on the phase diagram and composition. It is essential to develop novel Ti-based filler metals for optimizing the filler composition, regulating 3

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the microstructure, and improving the joint strength. To fabricate ribbon Ti-Cu-based filler metals, the glass-forming ability of alloy should be considered. In this study, Ti-Cu-based amorphous filler metals were designed based on the cluster-plus-glue-atom model put forwarded by Dong and Dong [18]. Binary Ti-Cu alloy with high glass-forming ability was selected as the matrix of Ti-Cu-based filler metals, aiming to obtain low-melting-point fillers. Ti and

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Cu are considered similar to Zr and Ni, respectively, due to weak ΔHTi-Zr=0 kJ/mol

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and ΔHCu-Ni=4 kJ/mol, where ΔH is the mixing enthalpy between elements [19, 20].

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Then partial Ti atoms could be substituted by Zr atoms, and so were Cu atoms by Ni

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atoms in Ti-Cu alloy. The resultant filler metal was expressed as Ti33.3Zr16.7Cu50-xNix

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(x=0-16.5%, atomic percent). The wettability of various filler metals on the substrates

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was contrastively discussed, and the effect of Ni content on the interfacial microstructure and shear strength of brazed joint was investigated. Furthermore, we

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characterized the reaction phases around the transition zone by EBSD and TEM, and revealed the crack initiation mechanism of the brazed joint. 2 Experimental procedure 2.1 Materials The raw materials of Ti, Zr, Cu and Ni with the purity of 99.99% were arc remelted under a vacuum of 6×10-3 Pa to obtain the homogeneous alloy ingots. Afterward, ribbon filler metals were fabricated with a melt-spinning method [15]. The alloy ingots were cut into small pieces and then put into the quartz tube with oblate mouth (5×1 mm2). Before melt spinning, the rotation speed of copper roller, the 4

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melting current, and the distance between oblate mouth of quartz tube and copper roller was set as 3000 r/min, 7 A, and 1 mm, respectively. The resultant filler foils are 5 mm wide and 80 μm thick. SUS 316L stainless steel and Ti6Al4V alloy sheets were applied as base metals with the chemical composition of Fe-16.8Cr-10.4Ni-2.0Mo-0.8Mn-0.6Si (wt.%) and Ti-5.9Al-3.6V (wt.%), respectively. To prepare the brazed specimens, 316L stainless

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steel sheets and Ti6Al4V alloy were cut into the size of 15×10×4 mm3 and 5×5×5

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mm3 with wire electrical machining, respectively. The faying surfaces to be brazed

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were gradually ground with up to 400 grit SiC papers, and then ultrasonically cleaned

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in acetone for 20 min.

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2.2 Wetting experiment

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The wetting test of filler metals on the substrates was carried out according to Chinese National Standards GB/T 11364-2008 [21]. Two substrates with the

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dimension of 40×40×2 mm3 were cut as the test specimens. Filler pellets weighing 0.1 g, cut from the alloy ingots, were used for the wetting test. The heating process for wetting experiment was consistent with brazing process. The holding temperature and time were set as 960 oC and 60 s, respectively. To investigate the influence of Ni content on the spreadability of filler metals, the drop areas were measured with AutoCAD software. Three wetting samples were measured to average drop areas under different filler metals. 2.3 Brazing procedure The filler foils with different Ni content were cut into small pieces with the 5

Journal Pre-proof dimension of 5×5 mm2, and then inserted between two base metals, as shown in Fig. 1. All the assembled workpieces tightly contacted under a pressure of 20 kPa by graphite blocks. The brazing experiment was carried out in a vacuum brazing furnace (ZTF2-10, China) under the vacuum of less than 5×10-3 Pa. According to preliminary studies, the brazing parameters were set as 990 oC for 10 min. The heating process

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was the same as our previous study [15]. To prepare metallographic samples, the joint

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polished with 1.5 m diamond polishing paste.

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cross-sections were gradually ground with up to 5000 grit SiC sandpaper and then

Fig. 1 Schematic diagram of brazed specimen

2.4 Characterization methods

The crystalline structure and thermal behavior of the filler metal were examined by an X-ray diffraction device (XRD, Empyrean) with Cu radiation and differential thermal analysis (DTA, TGA/DSC 3+), respectively. For XRD test, the operating voltage and current were set as 40 kV and 40 mA, respectively. Elemental distribution and interfacial microstructure of brazed joint were investigated by electron probe micro-analyzer (EPMA, JXA-8530F Plus) with a focused beam size of 100 nm in 6

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diameter. Electron back-scattered diffraction (EBSD) was used to characterize the reaction phases. The sample for EBSD was achieved using argon ion milling technique (PECS II MODEL 685) at 7 kV/90 min and subsequent 5 kV/30 min with a tilt angle of 2o. Scanning transmission electron microscope (STEM, JEOL JEM-2100F) was applied to identify the reaction phases. The TEM sample around the

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transition zone of the joint was prepared with focused ion beam (FIB, Helios G4 UX): (1) Deposit platinum over sampling location to protect the sample during ion beam

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thinning process; (2) Dig around the sample and pick up the sample with a probe; (3)

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Fix the sample to a copper half grid and thin the sample by ion beam at different

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voltages and tilt angles over and over. The resultant sample was about 50 nm thick.

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Shear test of the brazed joint was conducted using a universal testing machine

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(DNS-100) with the test speed of 0.5 mm/min, and three shear samples were used to average the shear strength of the joints brazed with different filler metals. Fracture

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surfaces of the brazed joint were examined by XRD and scanning electron microscope (SEM, ZEISS-SUPRA55) equipped with an energy dispersive spectroscopy (EDS, Oxford Instruments).

3 Results and discussion 3.1 Characterization of filler metals XRD patterns and DTA curves of Ti33.3Zr16.7Cu50-xNix filler metals are displayed in Fig. 2. The filler foils exhibit amorphous structure with existence of broad diffraction peaks at 2θ=42o when the content of Ni is 0, 5.5 and 11 at.% in Fig. 2(a). A few sharp crystalline peaks appeared in the XRD pattern of the filler foil with 16.5 at.% 7

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Ni, whereas the crystal structure of this filler metal matrix was principally amorphous. By calculating the inflection points around the endothermic peaks in Fig. 2(b), the melting temperature (Tm) and liquidus temperature (Tl) of the filler metals could be obtained, as listed in Table 1. The Tl of the filler metals with the Ni content of 0, 5.5 and 11 at.% ranged between 863 oC and 870 oC, and it reached 890 oC at 16.5 at.% Ni.

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Also, an endothermic peak appeared at about 950 oC when the Ni content was16.5 at.% as shown as the dotted circle marked in Fig. 2(b), due to the precipitation of

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crystalline phases CuTi in Fig. 2(a).

Fig. 2 (a) XRD patterns and (b) DTA curves of Ti33.3Zr16.7Cu50-xNix brazing filler metals Table 1 Melting temperature (Tm) and liquidus temperature (Tl) of Ti33.3Zr16.7Cu50-xNix filler metals Filler metals (at.%)

Tm (oC)

Tl (oC)

Tl-Tm (oC)

Ti33.3Zr16.7Cu50 Ti33.3Zr16.7Cu44.5Ni5.5 Ti33.3Zr16.7Cu39Ni11 Ti33.3Zr16.7Cu33.5Ni16.5

839 831 837 857

866 863 870 890

27 32 32 33

3.2 Wettability of filler metals on substrates The cross-section and top views of the wetting specimens were captured to calculate the spreading areas and contact angles of the filler metals on substrates, as displayed in Fig. 3. It can be seen that the filler metals have sufficiently spread on two substrates. For the Ti6Al4V wetting specimens, with increasing the Ni content from 0 8

Journal Pre-proof to 5.5 at.%, the contact angle slightly reduced from 13.5o to 12.3o at 16.5 at.% Ni, then it increased to a maximum value of 24.7o at 11 at.% Ni and finally dropped to 23.4o. For the 316L SS wetting specimens, with increasing the Ni content from 0 to 16.5 at.%, the contact angle first increased and then decreased with the peak value of 32.8o at 11 at.% Ni. All Ti6Al4V wetting specimens presented obvious reaction depth,

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which reached the maximum value at 11 at.% Ni, indicating the strongest reaction ability between this filler metal and the Ti6Al4V substrate. However, the 316L SS

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the droplet and 316L SS substrate.

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wetting specimens mainly exhibited flow spreading with a straight interface between

Fig. 3 Cross-section macro morphology wetting specimens with different filler metals: (a) and (b) Ti33.3Zr16.7Cu50; (c) and (d) Ti33.3Zr16.7Cu44.5Ni5.5; (e) and (f) Ti33.3Zr16.7Cu39Ni11; (g) and (h) Ti33.3Zr16.7Cu33.5Ni16.5

Fe-O oxide film limited the filler metal to spread on 316L stainless steel, however, Ti-O oxide film was not a limiting factor for filler metal spreading on Ti6Al4V alloy [22]. Furthermore, Ti-Cu-based filler metal has better compatibility with Ti6Al4V alloy than 316L stainless steel, which is favorable for reducing the interfacial tension between liquid filler metal and Ti6Al4V alloy, and thus the wettability was improved. The drop areas on Ti6Al4V substrates were larger than those on 316L SS substrates at the same Ni content, as shown in Fig. 4. When using 9

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Ti33.3Zr16.7Cu39Ni11 filler metal, the spreading areas on Ti6Al4V and 316L SS substrates reached the minimum of 40 mm2 and 30 mm2, respectively. Ni addition deviated the composition of the filler metal from the composition point with the superior glass forming ability, further leading to the improvement of melting temperature range (Tl-Tm), as shown in Table 1. At a larger melting temperature range,

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more solid phase(s) stayed in the filler metal, which reduced the wettability of the filler metal. Among all filler metals, Ti33.3Zr16.7Cu39Ni11 reacted with Ti6Al4V alloy

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most strongly, leading to the smaller but deeper spreading area.

Fig. 4 Drop areas of Ti6Al4V and 316L SS wetting specimen with various filler metals

3.3 Microstructure of brazed joints The brazed joints predominantly present two kinds of representative interfacial morphology, as exhibited in Fig. 5. One is for the brazed joints when Ni content of fillers was 0, 5.5 at.% and 11 at.%. The other is that with Ti33.3Zr16.7Cu33.5Ni16.5 filler metal. All brazed joints could be divided into three distinct zones (zones I, II and III). Same to our previous researches [15-17], zone I was a diffusion-reaction zone, 10

Journal Pre-proof including acicular (α+β)-Ti structure (Widmanstätten) adjacent to Ti6Al4V substrate and homogeneous diffusion zone (location A) close to zone II. When the Ni content in filler metals reached 16.5 at.%, β-Ti zone dissolved and migrated into zone II due to the precipitation of low-melting-point phase(s). Zone II gradually thickened with the increase of Ni content from 0 to 11 at.% and then intensively increased to the

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maximum thickness at 16.5 at.% Ni. Zone II consisted of the white phase (location B) and grey phase (location C). The grey phase reaction layer presented structural

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fluctuation, indicating that this layer was derived from the solidification with drastic

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reaction layers formed in zone III.

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energy fluctuation. In the high-magnification views shown in Fig. 5(a1)-(d1), three

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Fig. 5 Interfacial microstructure of the brazed joints for the filler metals with different content of Ni element: (a) and (a1) 0 at.%; (b) and (b1) 5.5 at.%; (c) and (c1) 11 at.%; (d) and (d1) 16.5 at.%

When the Ni content ranged from 0 to 11 at.%, the joint displayed the similar interfacial microstructure. The chemical composition of representative reaction phases in Fig. 5 is listed in Table 2. In zone I, spot A with a major constituent of 72.7 at.% Ti had the same phase contrast with intergranular phase of Ti6Al4V substrate. Also, 7.6 at.% Cu was identified. Cu element was a β-Ti stabilizer [23], which could induce the formation of β-Ti. In zone II, location B was retained from the filler metal because 12

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high content of Zr (15.2 at.%) was only found. Besides, Fe (10.6 at.%) and Cr (1.6 at.%) from the 316L SS substrate were detected, as well as Al (7.4 at.%) and V (3.4 at.%) from the Ti6Al4V substrate. The phase at location C was principally composed of 51.2% Ti + 14.0% Cu + 18.5% Fe (at. %), which was deduced to be (β-Ti + FeTi) according to Cu-Fe-Ti ternary phase diagram [24]. Table 2 EPMA chemical analysis results (at.%) of the marked positions in Fig. 5 Ti

Cu

Zr

Fe

Al

V

Cr

Ni

Possible phases

A B C D E F G H J

72.7 41.0 51.2 32.4 4.0 2.1 62.5 52.1 50.6

7.6 12.8 14.0 0.4 0.1 0.4 3.7 9.2 9.1

1.8 15.2 2.0 1.0 3.2 1.5 1.4

4.9 10.6 18.5 51.4 58.4 69.1 10.3 16.7 18.0

7.4 8.4 4.1 0.8 0.2 0.4 2.1 3.8 3.8

3.4 2.4 1.3 0.6 1.2 0.7 0.8 1.1 1.2

1.6 3.3 2.9 10.1 32.6 24.1 1.0 2.5 3.1

0.7 6.3 6.0 3.3 3.4 3.0 16.4 13.1 12.8

β-Ti residual filler metal β-Ti + FeTi (Fe, Cr)2Ti σ Fe-Cr β-Ti + NiTi2 (Cu, Ni)Ti2 + FeTi (Cu, Ni)Ti2 + FeTi

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Spots

In Fig. 6, Cr-rich and Ni-poor phenomenon could be seen in the transition zone

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III, which consisted of three reaction layers. At spot D, the content ratio of (Fe + Cr)

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(51.4 at.% + 10.1 at.%) and Ti (32.4 at.%) was approximately 2:1, which was correspondingly determined to be (Fe, Cr)2Ti according to Cr-Fe-Ti ternary phase diagram [25]. The reaction phases at spots E and F were σ phase (FeCr) and Fe-Cr solid solution, respectively. The σ phase was a binary intermetallic phase and detrimental to the brazed joints. It could be seen from Fig. 7 that Cr content in FeCr and Fe-Cr reaction layers was higher than that in both (Fe, Cr)2Ti layer and 316L SS substrate, likely attributed to the up-hill diffusion of Cr as discussed below.

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Fig. 6 Distribution of major elements near the transition zone corresponding to Fig. 5(d1)

Fig. 7 Line scanning results across the transition zone

When the Ni content of filler metal increased to 16.5 at.Fig%, two new reaction phases (spots G and H) formed. The phase at spot G was mainly composed of 62.5 at.% Ni, 10.3 at.% Fe, and 16.4 at.% Ni, which was determined to be (β-Ti + NiTi2) according to Fe-Ni-Ti ternary phase diagram [26]. Compared with spot G, the phase at location H had a lower content of Ti (52.1 at.%) and Ni (13.1 at.%), while higher content of Fe (52.1 at.%), which was consistent with the mapping distribution of Ti, Fe, and Ni in Fig. 6. Furthermore, 9.2 at.% Cu was detected, which was regarded as a 14

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similar element to Ni. So the phase at spot H was characterized as (Cu, Ni)Ti2 + FeTi, same to the phase at location J in zone II adjacent to zone III. EBSD was used to further investigate the reaction phases near zone III for the typical joint brazed with Ti33.3Zr16.7Cu39Ni11 filler metal, as shown in Fig. 8. The crystal structure of residual filler metal in zone II was distorted with various major components (Ti, Zr, Cu, Fe, and Ni), resulting in low EBSD resolution in this zone.

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The reaction phase of (β-Ti + FeTi) close to zone II was principally identified to be

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FeTi, and the (Fe, Cr)2Ti phase in zone III was mainly regarded as Fe2Ti phase with

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lower resolution. The recognition for FeCr (σ) polycrystal was sufficient in the middle

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of zone III. However, Fe-Cr phase close to 316L SS substrate (γ-Fe) was completely

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labeled as FeTi phase, which did not match the above analysis on the microstructure

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in Fig. 5. The reason for mislabeling was that Fe-Cr phase likely had an extremely similar crystal structure with FeTi, which is body-centered cubic (BCC) with

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a=b=c=0.2976 nm. In the Fe-Cr system, only FeCr (σ), γ-Fe and α-Fe appeared, and α-Fe had the similar BCC crystal structure and lattice constants to FeTi phase with a=b=c=0.29315 nm [27]. Thus, Fe-Cr phase was deduced to be α-Fe. Generally speaking, the EBSD resolution was not ideal in this study, however, it contributed to the identification of the reaction phases to some degree.

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Fig. 8 EBSD image of reaction phases

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To confirm the reaction phases within the joint brazed with Ti33.3Zr16.7Cu39Ni11

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filler metal, the investigation with TEM was conducted, as shown in Fig. 9. FIB

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sampling position was mainly focused on near the transition zone (zone III) in Fig,

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9(a). The TEM sample in Fig. 9(a1) before milling by ion beam was compared with

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the TEM bright-field view to confirm the locations of different reaction layers. In Fig. 8(b), the TEM bright-field view can be accordingly divided into four distinct zones,

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and relevant selected area electron diffraction (SAED) patterns are shown in Fig.

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9(c)-(f). The patterns from SAED 1 to SAED 4 were acquired along the zone axes of [111], [112], [100], and [111], respectively, and each SAED pattern revealed the single phase. SAED 1 confirmed FeTi phase without the detection of β-Ti. However, 14 at.% Cu (spot C, Table 2) was detected in this zone, which was a strong β-Ti stabilizer. Upon solidification, Ti preferred to react with Fe due to their larger enthalpy of mixing (ΔHTi-Fe=-17 kJ/mol) than Ti-Cu (ΔHTi-Cu=-9 kJ/mol) [20]. Then surplus Ti would transfer into β-Ti phase with extremely similar crystal structure to FeTi phase, as discussed above. So SAED 1 was confirmed to be (β-Ti + FeTi). SAEDs 2 and 3 demonstrated Fe2Ti and FeCr phases, respectively, which was consistent with the EBSD analysis results. SAED 4 in Fig. 9(f) revealed the α-Fe phase at location F in 16

Journal Pre-proof Fig. 9. So, up-hill diffusion of Cr would promote the γ-Fe to transfer into α-Fe with ~30 at.% Cr. So the interfacial microstructure of the brazed joint was finally expressed as titanium substrate/Widmanstätten/β-Ti/residual filler metal/(β-Ti +

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FeTi)/Fe2Ti/FeCr/α-Fe/316L SS substrate (γ-Fe).

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Fig. 9 TEM analysis results: (a) FIB sampling position for TEM sample; (b) TEM bright-field image cross the interface; (c)-(f) selected-area electron diffraction patterns corresponding to the marked locations in (b)

According to Fe-Ti binary phase diagram [28], the melting points of FeTi in zone II and Fe2Ti in zone III are 1317 oC and 1427 oC, respectively, higher than the brazing temperature 990 oC. And (Fe, Cr) solid solution transforms into solid FeCr phase below 830 oC and liquid phase above 1513 oC with about 21 at.% Cr, respectively. During brazing process, once reaching the liquidus temperature, the filler metal began to melt, and two solid/liquid interfaces formed between the molten filler metal and base metals. Through the molten filler metal/316L SS substrate interface, the mutual diffusion of Ti and Fe occurred, thus, the melting point of filler metal increased with Fe continually diffusing into the molten filler metal. At a certain Fe content, the 17

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molten filler metal began to solidify with the formation of FeTi in zone II adjacent to zone III. Simultaneously, Ti diffused into 316L SS substrate and Fe2Ti phase formed at a certain Ti content. Ti content was the highest in the Fe2Ti layer of zone III due to short-range diffusion. Then Ti further diffused toward 316L SS substrate with long-range diffusion, forming (Fe, Cr) solid solution with different ratios of Fe to Cr in different zones. Upon cooling, (Fe, Cr) solid solution transformed into FeCr layer

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below 830 oC and α-Fe layer, respectively. Diffusion of Ti into 316L substrate

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weakened the Cr activity and reduced Cr solubility in 316L SS substrate, resulting in

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up-hill diffusion and segregation of Cr in FeCr and α-Fe reaction layers. This

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phenomenon was also confirmed in Refs [29, 30].

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3.4 Mechanical properties

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The compression-shear test was conducted to investigate the effect of Ni level in filler metals on the joint shear strength. To exactly evaluate the shear strength of the

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brazed joint, a specific fixture for shear test was machined, as illustrated on the top right in Fig. 10(a). The brazed joint was inserted in the fixture, and the position of brazed seam was vertical. The shear load was applied on the titanium alloy substrate with a vertical pressure until the joint failure. The shear strength was 163 MPa at 0 Ni, then increased to 220 MPa with 5.5 at.% Ni filler metal. Increasing the Ni content to 11 at.%, the brazed joint reached the maximum shear strength of 318 MPa. While further going to 16.5 at.% Ni, the joint shear strength sharply decreased to 114 MPa. As discussed on the wettability, increasing the Ni content modified the viscosity and flowability of filler metal. When the Ni content was lower (0 and 5.5 at.%), the 18

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molten filler metal was easy to flow out of the brazed seam, leading to insufficient diffusion and reaction between filler metal and substrates. However, at higher Ni content of 16.5 at.%, most of the molten filler metal stayed in the seam and excessively reacted with the substrates, resulting to the formation of abundant brittle compounds and deteriorating the joint properties. Adding 11 at.% Ni in the filler metal

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ensured the filler metal with appropriate viscosity and flowability for brazing Ti6Al4V alloy/316L SS. The diffusion and reaction between filler metal and

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substrates were sufficient for metallurgical bonding in Ti6Al4V alloy/316L SS joint,

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leading to high joint strength. For comparison, Fig. 10(b) illustrates other researchers’

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results on the strength of brazed titanium alloy/steel joint, including various brazing

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couples, filler metals, and strength types [9, 10, 31-38]. All shear strength in other

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studies is smaller than that in this study, and only tensile strength in Ref. [10, 37, 38] was larger than the shear strength in this study, which demonstrates the advantage of

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Ti33.3Zr16.7Cu39Ni11 amorphous filler metal on brazing of Ti6Al4V alloy/316L SS.

Fig. 10 Joint strength: (a) Influence of Ni content in filler metals on the shear strength of brazed joints in this study; (b) Comparison of joint strength between this study and other studies

3.5 Fracture analysis The fracture paths of brazed joints after shear test are shown in Fig. 11. All 19

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brazed joints fractured along the interface of zone II/zone III in Fig. 11(a)-(d). From the high-magnification views in Fig. 11(a1)-(c1) and Fig. 11(a2)-(c2) with 0-11 at.% Ni filler metals, it could be further seen that the cracks initiated at the interface of (β-Ti + FeTi)/Fe2Ti, and propagated along the reaction layers of Fe2Ti and FeCr. When the Ni content of filler metal was 11 at.%, a large amount of rough secondary cracks

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appeared in Fig. 11(c)-(c1), which could absorb more shear energy. When the Ni content was 16.5 at.%, the crack propagation path changed, compared with 0, 5.5, 11

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at.% Ni filler metals. Similarly, the cracks initiated at the interface of zone II/zone III,

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essentially the interface of reaction layers ((Cu, Ni)Ti2 + FeTi)/Fe2Ti. Afterward, the

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cracks mainly propagated within the layers of (Cu, Ni)Ti2 + FeTi (zone II), Fe2Ti

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(zone III) and FeCr phase (zone III). As discussed above, the interfaces of (β-Ti +

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FeTi)/Fe2Ti and ((Cu, Ni)Ti2 + FeTi)/Fe2Ti were solid/liquid interfaces during solidification, which were the weak part of brazed joints with large stress

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concentration. In addition, FeTi, Fe2Ti, (Cu, Ni)Ti2 and FeCr phases were brittle intermetallic compounds, and their lattice mismatch would severely deteriorate the brazed joints.

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Fig. 11 Cross-section views of fracture path for the failed joint brazed with different Ni content filler metals (a) 0Ni, (b) 5.5Ni, (c) 11Ni, (d) 16.5Ni; (a1)-(d1) high-magnification images of Ti6Al4V alloy side; (a2)-(d2) high-magnification images of 316L SS side

The fracture surfaces of the failed joints brazed with Ti33.3Zr16.7Cu39Ni11 and Ti33.3Zr16.7Cu33.5Ni16.5 filler metals are displayed in Fig. 12, and the chemical composition of particular positions is listed in Table 3. On both Ti6Al4V alloy and 316L SS sides, β-Ti + FeTi (locations A and G), Fe2Ti (locations B and C), and FeCr phase (location D and E) were detected. The above reaction phases possessed specific fracture features. In Fig. 12(e), the (β-Ti + FeTi) phase mainly presented the cleavage feature with the presence of massive cleavage steps and facets. The fracture surface, 21

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where Fe2Ti located in, was a bit flat with lots of tiny folds, same to that in Fig. 12(f). The lattice mismatch between cubic FeTi and hexagonal Fe2Ti phases might induce the intergranular fracture with the flat surface of Fe2Ti phase. FeCr phase revealed intergranular fracture characteristic in Fig. 12(c). For the failed joint brazed with Ti33.3Zr16.7Cu33.5Ni16.5 filler metal, a large number of smooth planes and cleavage steps

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were observed on the Ti6Al4V alloy side in Fig. 12(g), indicating brittle fracture feature. The cleavage fracture is a kind of brittle transgranular fracture, which easily

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occurs on some particular crystal planes with low index [39]. On the 316L SS side,

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layered fracture morphology could be observed. In Fig. 12(g) and (f), (Cu, Ni)Ti2 +

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FeTi phases (zone II in Fig. 5(d)) mainly displays cleavage (Fig. 12(d)) mixed with

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intergranular features with the existence of smooth planes.

Fig. 12 Fracture surfaces of the joints brazed with (a) and (b) Ti33.3Zr16.7Cu39Ni11; (g) and (h) Ti33.3Zr16.7Cu33.5Ni16.5; (c)-(f) corresponding high-magnification views 22

Journal Pre-proof Table 3 EDS chemical analysis results (at.%) of the marked locations in Fig. 12 Ti

Cu

Zr

Fe

Al

V

Cr

Ni

Possible phases

A B C D E F G H I J

46.6 32.7 28.4 7.1 5.7 50.3 47.3 30.1 3.5 43.9

19.4 1.0 0.6 0.2 12.6 17.6 0.3 0.3 17.8

1.5 2.3 4.8 0.5 0.4 1.8 1.9 1.8 0.4 1.8

19.6 51.5 47.2 58.9 56.5 15.8 19.1 52.9 58.8 16.1

3.6 0.6 1.1 0.7 0.4 4.6 5.1 0.7 0.3 5.6

0.7 0.5 1.2 1.0 1.3 1.6 1.6 0.5 1.4 1.4

1.9 7.1 13.3 28.8 32.5 2.8 2.4 9.7 31.5 3.1

6.8 4.4 3.4 3.2 3.0 10.5 5.0 4.2 3.9 10.3

β-Ti + FeTi Fe2Ti Fe2Ti FeCr FeCr (Cu, Ni)Ti2 + FeTi β-Ti + FeTi Fe2Ti FeCr (Cu, Ni)Ti2 + FeTi

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Figure 13 exhibits XRD patterns of the fracture surfaces brazed with

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Ti33.3Zr16.7Cu39Ni11 and Ti33.3Zr16.7Cu33.5Ni16.5 filler metals. While using

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Ti33.3Zr16.7Cu39Ni11 filler metal, FeCr (σ), Fe0.8Cu0.2Ti and Fe2Ti reaction phases were

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verified on both fracture surfaces of Ti6Al4V alloy and 316L SS sides. Fe0.8Cu0.2Ti was regarded as FeTi phase (zone II, Fig. 5(c1)) with the same space group of Pm-3m

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and similar cell parameters (a=b=c≈0.3 nm, α=β=γ=90o). However, while using

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Ti33.3Zr16.7Cu33.5Ni16.5 filler metal, Fe0.2Ni4.8Ti phase, which had the same space group and similar lattice with Fe0.8Cu0.2Ti and FeTi, was detected on the fracture surfaces. So Fe0.2Ni4.8Ti was also deemed to be FeTi phase, corresponding to zone II in Fig. 5(d1). In addition, Ti substrate and Fe substrate were also found on respective fracture surfaces. The XRD detection on the fracture surfaces was fully supportive to the above analysis on the discussion about fracture paths.

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3.6 HRTEM of FeTi/Fe2Ti interface

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Fig. 13 XRD patterns of the fracture surfaces

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To interpret the fracture mechanism of the brazed joint, the high-resolution TEM

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(HRTEM) image near the FeTi/Fe2Ti interface was captured, as shown in Fig. 14. According to the fast Fourier transformation (FFT) in Fig. 14(c) and (d), the HRTEM

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images of FeTi and Fe2Ti phases were obtained along the [111] and [112] zone axes, respectively. The crystallographic orientations of both FeTi and Fe2Ti in HRTEM -

image were (011). In Fig. 14(b), the FeTi/Fe2Ti interface was chaotic, which was likely attributed to the overlapping crystal planes due to the tilted zone axes of FeTi and Fe2Ti phases. To reduce the interfacial energy, two different phases were demanded with low lattice mismatch, less than 6%, thus obtaining high joint strength [40]. The lattice mismatch of FeTi/Fe2Ti interface could be calculated by the following equation [41]: F=2(dFe2Ti -dFeTi )/(dFe2Ti +dFeTi ) 24

Journal Pre-proof where dFe2Ti (0.398 nm) and dFeTi (0.211 nm) are d-spacings of Fe2Ti and FeTi, respectively, as shown in Fig. 14(b). Thus, the lattice mismatch F between Fe2Ti and FeTi was calculated to be 61.4%, which indicated that the FeTi/Fe2Ti interface belonged to a non-coherent interface with high strain energy. Plenty of dislocations -

would form to reduce the strain energy. During shear test, the slip of (011)

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crystallographic plane was restrained. Once the shear stress exceeded the strain energy, the brazed joint would initiate to fracture along the FeTi/Fe2Ti interface with brittle

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fracture mode.

Fig. 14 (a) TEM bright-field image of the FeTi/Fe2Ti interface; (b) HRTEM image near the FeTi/Fe2Ti interface; (c) and (d) FFT patterns of FeTi and Fe2Ti, respectively

4 Conclusions A series of Ti-Cu-based amorphous filler metals (Ti33.3Zr16.7Cu50-xNix) were 25

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designed to evaluate the effect of Ni level on the wettability of filler metals, interfacial microstructure and shear strength of Ti6Al4V alloy/316L stainless steel brazed joints. The main conclusions are listed as follows. (1) Ni addition deviated the composition of the filler metal from the eutectic point of Ti-Cu, reducing the wettability of filler metals. The filler metal with 11 at.%

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Ni possessed inferior wettability but was optimized for brazing of Ti6Al4V alloy/316L stainless steel with the maximum joint shear strength of 318 MPa.

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(2) The brazed joints presented a particular layered interfacial microstructure of

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titanium substrate/Widmanstätten/β-Ti/residual filler metal/(β-Ti +

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FeTi)/Fe2Ti/FeCr/α-Fe/316L stainless steel substrate (γ-Fe). The reaction phases,

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including (β-Ti + FeTi), Fe2Ti, FeCr and α-Fe, generated near the transition zone. Ni

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addition in filler metals thickened the brazed seam due to the weakening wettability. (3) Lattice mismatch of 61.4% between FeTi and Fe2Ti phases initiated cracks at

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the (β-Ti + FeTi)/Fe2Ti interface. The initiative cracks mainly propagated within the reaction layers of Fe2Ti and FeCr with brittle feature. Different reaction phases revealed peculiar fracture morphologies, including cleavage (β-Ti + FeTi), brittle tiny folds (Fe2Ti), and intergranular (FeCr). Acknowledgements This work was financially supported by the National Natural Science Foundation of China (51674060), the Fundamental Research Funds for the Central Universities (DUT18LAB01), and technically supported by the Collaborative Innovation Center of Major Machine Manufacturing in Liaoning. 26

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Journal Pre-proof Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

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Journal Pre-proof Author’s Statement Transparency about the contributions of authors as below:

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Yueqing Xia conducted the experiment and draft the manuscript. Honggang Dong guided the experiments and finalized the manuscript. Runze Zhang helped braze the joints. Yaqiang Wang helped SEM experiments. Xiaohu Hao assisted in EPMA experiments. Peng Li helped TEM experiments. Chuang Dong helped design the filler metals.

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Graphical abstract

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Highlights Ti33.3Zr16.7Cu50-xNix (x=0 to 16.5%) amorphous filler metals were designed.

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The maximum joint shear strength was 318 MPa, higher than reported results.

3.

The reaction phases around the transition zone were confirmed.

4.

The lattice mismatch between FeTi and Fe2Ti phases was 61.4%.

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