RM80 steel dissimilar resistance spot welding joint

Materials Science & Engineering A 776 (2020) 139008

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Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea

Microstructure and mechanical property of WC-10Co/RM80 steel dissimilar resistance spot welding joint Gang Chen a, Wei Xue a, *, Yuzhen Jia b, c, Shucheng Shen a, Guoyue Liu b, c a

College of Materials Science and Engineering, Hunan University, Changsha, 410006, China Sawing Engineering Research Center of Hunan Province, Changsha, 410200, China c Bichamp Cutting Technology (Hunan) Co, Ltd., Changsha, 410200, China b

A R T I C L E I N F O

A B S T R A C T

Keywords: Dissimilar resistance spot welding RM80 steel WC-10Co η phase Mechanical property High speed camera

In this work, resistance spot welding (RSW) was used to weld dissimilar metals of cemented carbide (WC–10Co) and high strength steel (RM80). The mechanical property and microstructure of welded joints were analyzed by shear test, scanning electron microscopy and micro X-ray diffraction, establishing the corresponding relationship between the mechanical property of welded joints and microstructure. The results show that with the increase of welding current, the shear strength of welded joints first increased and then decreased, the maximum of which under the optimal welding process was 924 MPa. The characteristics of five interface bonding types were clarified, whose proportion under different welding currents was calculated, and the contribution of each interface bonding type to the performance of welded joints was studied based on the mechanical property and fracture mechanism.

1. Introduction Cemented carbide, due to its high hardness, good wear resistance and corrosion resistance [1–4], is widely applied in aerospace, machining, oil drilling, mining tools and other wear-resistant compo­ nents [5–7]. More than 60% of cemented carbide is used to make cutting tools, such as cutter heads, drill bits, band saw blade, etc. In order to reduce the influence of intrinsic brittleness as well as control cost, cemented carbide is often used as the tooth embedded into high strength and high toughness steel, forming high strength and high toughness cutting tools and drills with high hardness and wear resistance tooth [8]. In particular, the cemented carbide band saw blade for sawing high-alloy steel, titanium alloy, stainless steel, superalloy and other difficult-to-machining materials is formed by welding WC-10Co cemented carbide and RM80 steel. It can be seen that the key to the preparation of cemented carbide drilling tools lies in its weldability with steel. The quality of welding undoubtedly determines the working efficiency, service life and cost of the tool. Therefore, the welding of cemented carbide and steel dissimilar metals was widely researched for a long time, with the focus on the bonding strength of welded joints. For example, using different welding processes [9–13], adopting composite interlayer [13–17] and

introducing weld preheating or post-weld heat treatment [18] can improve the welding strength. The large difference in thermal expansion coefficient (TEC) between cemented carbide and steel [19] and large wetting angle of cemented carbide on steel, along with the brittle phase of weld fusion zone is the main influential factors for strength, which are also challenges facing the welding of cemented carbide and steel dis­ similar metals. In fact, the widely used active metal (or suitable metal such as Ni) interlayer, along with weld preheating and post-weld heat treatment, have solved the problems including poor wettability and large differences in TEC when welding cemented carbide and steel. During the welding process of cemented carbide and steel dissimilar metals, the η phase will be formed due to the diffusion of C element in this zone [20–22]. Owing to the high bulk modulus of η phase [23], which is much larger than the weld material, interface micro-cracks will occur after welding and the strength of welded joints will be seriously affected [12,24,25]. Massive η phase can be found in welded joints in brazing, electron beam welding, laser welding and diffusion welding. Increasing the brazing current [10], brazing time [26], or introducing ultrasonic [27] all contributes to the formation of the η phase. Increasing laser welding energy density [21], or the bonding time of diffusion welding [28], or other process parameters [9,29], also increases the content of η phase in the welded joint. However, when friction stir

* Corresponding author. E-mail address: [email protected] (W. Xue). https://doi.org/10.1016/j.msea.2020.139008 Received 29 November 2019; Received in revised form 20 January 2020; Accepted 23 January 2020 Available online 28 January 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.

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Fig. 1. (a) Schematic diagram for RSW, and (b) as-welded sample.

promote the formation of carbide solid solution phase (CSSP) which is adjacent to the η phase [29], thereby suppressing the formation of the η phase, due to the fact that CSSP has smaller Gibbs free energy than the η phase. The number and distribution of the η phase are also related to trace elements in cemented carbide and surface chemical composition. For example, adding 0.5% Cr3C2 into WC-8Co cemented carbide [33] and WC-Co (45 wt%.-WC~85 wt%.-WC) cemented carbide subject to surface nitriding treatment [18] can also inhibit the formation of η phase at the weld zone. Therefore, the brittle phase (η phase) at the interface of welded joints in welding cemented carbide and steel dissimilar metal is a core problem that needs to be solved urgently. This study is based on the preparation of cemented carbide band saw blades. The RSW technology, which has good metallurgical bonding compared with brazing, ensures high bonding strength and overcomes the limitations of high energy beam welding such as laser welding for small-sized joints, is used for the welding of WC-10Co cemented carbide and RM80 steel dissimilar metal. The study explores welding currents (I) and microstructure (interface bonding type), combines strength and fracture mode, and obtains excellent mechanical property under suitable process conditions.

Table 1a Chemical compositions of RM80 steel (wt.%). RM80

C

Si

Mn

Cr

Ni

Mo

V

Fe

0.38

0.39

0.69

2.97

0.65

2.11

0.28

Bal.

Table 1b Compositions of WC-10Co cemented carbide (wt.%). WC-10Co

WC

Co

Cr2C3

89.5

10

0.5

welding was adopted for welding cemented carbide and steel dissimilar metals, as the thermal cycling only lasts a few seconds, reducing the diffusion of Fe, Co and C elements can significantly decrease the number of η phase at the weld interface [30]. Similarly, when solid capacitor spark sintering welding is utilized for welding WC-12Co and steel, as the short discharge time lasts only 20 ms, the metallurgical reaction be­ comes insufficient, so the η phase is not revealed at the welding interface [31]. In addition, presetting Ni interlayer between cemented carbide and steel can also significantly reduce the η phase in the weld zone [8, 32], because elements such as Ni, Fe and Co can dissolve mutually and form a solid solution which can effectively prevent the formation of the η phase. In addition, adding 0.5 wt% La2O3 to the Cu interlayer will

2. Material and methods 2.1. Material This study is based on cemented carbide band saw blade, using RM80 steel as the backing material, with the size of 1.3 mm � 41 mm and chemical composition shown in Table 1a. The steel band is milled to form a tooth shape (Fig. 1a), and austenitized at 1150 � C and tempered at 600 � C to obtain a martensite microstructure and a hardness of 500 HV0.3; the WC-10Co cemented carbide cylinder with the dimension of Φ1.98 mm � 2.40 mm will be used as the tooth of band saw blade. Chemical composition shown in Table 1b, and the hardness is 1650 HV0.3.

Table 2 Process parameters for RSW of WC-10Co/RM80 steel. Sample 1 Sample 2 Sample 3

Preheat current

Preheat time

Welding current

Welding time

300A 300A 300A

20 ms 20 ms 20 ms

650A 800A 950A

95 ms 95 ms 95 ms

Table 3 EDS analysis results of the WC-10Co/RM 80 steel interface marked in Fig. 6. Regions

C

Fe

O

W

Co

Ni

Fe/W(wt.%)

Fe/W(at.%)

A B C D E

25.48 20.11 23.60 25.28 16.73

46.65 31.86 62.24 49.86 69.61

13.69 32.04 7.94 8.6 8.21

8.90 15.56 3.63 8.15 1.00

4.46 0.43 2.20 5.22 3.82

0.82 0.00 0.39 2.89 0.63

5.24 2.05 17.15 6.12 69.61

17.21 6.73 56.34 20.10 228.67

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Fig. 2. (a) Schematic and (b) picture of device for shear test.

2.2. RSW process

10Co/RM80 joint interface and the fracture surface. Micro X-ray diffraction (micro XRD) was used to detect the phase of weld and frac­ ture surface (steel side and cemented carbide side). Images of the RSW process were recorded by a FASTEC Hispec5 high speed camera at the frame rate of 1708 fps. The temperature changes of the weld zone under different currents were described by the IGAR12-LO colorimeter py­ rometer. The detection area of the colorimeter pyrometer was a circular spot with a diameter of 1 mm, where the average temperature was detected. The mechanical property was evaluated by shear test, whose sche­ matic and picture are shown in Fig. 2. The test was performed at an indenter speed of 0.5 mm/min at room temperature, which could be considered as a static test [34]. The shear force was read from a com­ puter connected to the universal tester sensor, and the shear strength value was obtained based on the calculation method provided in G.W. Liu and F. Valenza [35].

In this experiment, the welding of cemented carbide band saw tooth and the backing material adopted the saw-tooth micro-welding tech­ nology special machine. Before welding, the groove was preseted at the welding position of steel band, and the surface of cemented carbide cylinder was plated with a C-containing Ni-based interlayer with the thickness of 5–10 μm, which can relieve residual stress in the fusion zone and increase electrical conductivity. The RSW process parameters are shown in Table 2, and the schematic diagram and post-weld macro­ scopic morphology are shown in Fig. 1a and b. The post-weld sample was subjected to the stress relief annealing treatment at 400 � C/15min in a nitrogen atmosphere. 2.3. Microstructures and mechanical property characterization The etching of welding interface was carried out by using Murakami reagent (20 g K3[Fe(CN)]6, 20 g NaOH and 80 g H2O, freshly config­ ured). The distribution and morphology of WC and η phase were observed by OLYMPUS DSX510 optical microscope (OM). Microstruc­ tures and fracture morphology were revealed by FEI QUANTA200 scanning electron microscope (SEM) with energy dispersive X-ray (EDAX), which was used to determine the distribution of element of WC-

3. Results 3.1. Weld pool evolution The RSW pools evolution under different welding currents is shown in Fig. 3. Fig. 4 depicts that the temperature change of zone A (Fig. 4a,

Fig. 3. RSW pool evolution under different welding currents: (a) 650A, (b) 800A, and (c) 950A 3

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Fig. 4. The temperature change of zone A (Fig. 4a, d and g), zone B (Fig. 4b, e and h) and zone C (Fig. 4c, f and i) of the welded joint under 650A (Fig. 4a, b and c), 800A (Fig. 4d, e and f) and 950A (Fig. 4g, h and i). The detection position of the colorimetric pyrometer is shown in the red round spot of each figure. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

d and g), zone B (Fig. 4b, e and h) and zone C (Fig. 4c, f and i) of the welded joint under 650A (Fig. 4a, b and c), 800A (Fig. 4d, e and f) and 950A (Fig. 4g, h and i). The detection position of the colorimetric py­ rometer is shown in the red round spot of each figure. t ≥ 800� C, t ≥ 1300� C and t ≥ 1800� C are defined as the duration at the temperature above 800 � C, 1300 � C and 1800 � C, respectively. When the welding current was 650A, it was found that the middle area at the contact interface slightly softened and deformed at 15 ms. With the increase of time, the contact interface of WC-10Co cemented carbide and RM80 steel gradually increased during the welding process. The highest temperature was always reached at the center of the middle area, while the weld slag (partially melted cemented carbide and steel) and current density were revealed on the edges of contact interface. When it reached 115 ms, the temperature of the weld reached the maximum and caused severe deformation. As the welding current increased to 800A and 950A at 15 ms, more obvious deformation was found in the middle area and changes were experienced by the tem­ perature field, and the time was also advanced when the temperature of the weld and the deformation degree of the joint reached the maximum. With the welding current increased, both t ≥ 800� C and t ≥ 1300� C increased first and then decreased as shown in Fig. 4a, d and g. Although the t ≥ 800� C and t ≥ 1300� C of the weld center area (zone A) at the welding current 950A is shorter than those at the welding current 800A, the t ≥ 1800� C of the weld center area (zone A) under the welding current 950A is

about 10–20 ms, and the maximum temperature of this area did not reach 1800 � C at the welding current of 800A. In Fig. 4b, e and h, it can be clearly observed that the t ≥ 800� C of zone B also increased first and then decreased with the increase in welding current, which verifies the temperature change law in the center of the weld area (zone A). Compared with time, the maximum temperature has a far greater impact on the weld structure, which reflects the effective welding length increased with the increase of duration at high temperature. However, when the welding current was 950A, the center area of the weld heated up and cooled down much faster than at the welding currents of 650A and 800A. This also confirms that the excessive welding current in Figs. 3c and 5c made the melting area on the steel side too large, and the welding slag was extruded from the weld and contacted the electrode, causing the heat of the welding joint to be quickly transmitted through the electrode, and cracks were formed in the welding joint, which seriously affects the bonding strength of the welded joints. As it moves away from the central area of the weld, its temperature continues to drop. 3.2. Microstructure characterization Fig. 5 illustrates the OM images of the joint after RSW under different currents. The upper part is the tooth material of WC-10Co cemented carbide, and the lower part is RM80 steel. As the welding current 4

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welding current is 650A and 950A, while F and G patterns are obtained from the fracture surface of the RM80 steel side. The EDS analysis results of elemental content in each region are given in Table 3. In interface bonding type-A, only γ-Co dissolved in the WC-10Co cemented carbide, and the Fe were bonded by diffusion forming a solid solution, while the edges of the WC particles did not appear to be passivated. It can be seen from the EDS analysis of Fig. 6b that a small number of Fe elements entered the zone I/II of WC-10Co cemented carbide, while W element was not found on zone V, and the width of the zone III was only about 1 μm. In interface bonding type-B, the edges of WC particles in the zone III dissolved, the particles near the steel side are almost completely “spheroidized”, and a small number of η phase may be formed in the vicinity of the “spheroidized” WC particles. Compared with interface bonding type-A, the average width of fusion zone for typeB doubled, and the distribution of the C element content became rela­ tively moderate. It is worth noting that the WC melting process was also a metallurgical reaction process with elemental diffusion and grain growth. Compared with interface bonding type-A and type-B, type-C (Fig. 6e) clearly highlights the loose zone (zone II) and fusion line (zone IV). The zone II of WC-10Co cemented carbide was formed as massive Fe atoms in RM80 steel migrate to WC-10Co cemented carbide substrate at high temperature. The width of zone III for interface bonding type-C continues to grow to about 4 μm, and the width of the zone IV is longer than 1 μm. The microstructure of type-C was mainly “spheroi­ dized” WC particles, Fe and a few η phase, which proves that WC-10Co cemented carbide and steel carried out sufficient metallurgical re­ actions. It can be found in Fig. 6b, d and f that there is an oxygen peak in the fusion zone near the side of WC-10Co cemented carbide. As the width of fusion zone increased, the width of oxygen peak also increased, but the remain basically consistent in height. The fusion zone structure of type-D is completely different from other interface bonding types, in which massive fishbone or herringbone eutectic structures were formed, whose interior was austenite and exterior was η phase or WC phase [9]. The average width of zone III reached 10 μm, the largest compared with others. The width of zone IV was about 3 μm, which can reflect the degree of metallurgical reaction of WC-10Co cemented carbide and steel to some extent. The wider the zone III is, the more sufficient the reaction between the two becomes. It can be clearly seen from Fig. 6i that the zone III for type-E was extruded into the zone V by some external forces. The white portion in Fig. 6j is “spheroidized” WC particles, the black portion may be Fe, and the gray portion is η phase which is formed to be dispersed between Fe and “spheroidized” WC particles [13,22]. Fig. 6k exhibits the crack near the welded interface of WC-10Co cemented carbide/RM80 steel, which is indicated by the red arrow in Fig. 5c. Fig. 8 depicts the percentage of five interface bonding types at the weld under the welding currents 650A, 800A and 950A. The statistical model uses a method that replaces the volume with the length of interface bonding type along the weld direction (The more information is provided in the Supplemental Information: B). When the welding current was 650A, the type-A occupied a dominant position, exceeding 75%. As the welding current increased to 800A and 950A, the propor­ tion of type-A first decreased and then increased. The percentage of type-B, type-C, type-D and type-E first increased and then decreased with the increase of welding current. When the welding currents was 650A and 950A, the proportion of the type-A was the largest, while the type-B dominated at the welding current of 800A. It is worth noting that when the welding currents increased from 800A to 950A, the proportion of type-E dropped sharply. There is a possibility that due to the action of welding pressure, excessive thermal cycling forced the WC particles in the fusion zone to be extruded toward the sides of welded joint with the slag. It can be noted in Fig. 5c that the reddish brown areas on both sides are larger than those in Fig. 5b, which provides a strong evidence.

Fig. 5. OM images of RSW joints under different welding currents: (a) 650A, (b) 800A, and(c) 950A.

increased, the fusion area of the welded joint expanded, and the extruded slag also increased. It is worth noting that cracks occurred near welded joint (cemented carbide side) at the welding current 950A, as indicated by the red arrow in Fig. 5c, and the crack didn’t appear at the welding currents 650A and 800A, which is caused by the large differ­ ences in TEC between the two materials and the extremely rapid rise and drop of the temperature. Fig. 6 shows the SEM micrographs of five interface bonding types in the weld zone, the EDS results and the SEM micrograph of the weld crack; the welded joints have five types of interface bonding shown in Fig. 6a, c, e, g and i, which are respectively classified as type-A, type-B, type-C, type-D and type-E (Further details of the interface combined with the classification of types is provided in the Supplemental Infor­ mation: A Fig. B.3). Meanwhile, along the direction perpendicular to the weld seam, each interface bonding type can be divided into five areas according to the difference in structure, which are recorded as: zone IHAZ near the WC-10Co cemented carbide side (HAZ (WC–10Co)); zone II-loose zone; zone III- fusion zone; zone IV- fusion line; zone V- the HAZ of RM80 steel (HAZ (RM80)). Fig. 6j shows the partial enlarged view of the fusion zone for interface bonding type-E. Fig. 6k depicts the back­ scattered electron micrograph of the crack near the welding interface of WC-10Co cemented carbide/RM80 steel, as shown in the red arrow of Fig. 5c. Fig. 6b, d, f and h show the distribution of W, Fe, C, O and Co elements from type-A to type-D in the welds perpendicular to the weld direction. Various phases are identified by the micro XRD on weld and fracture surface, as shown in Fig. 7, where A pattern is obtained from the position of interface bonding type-D, B and C patterns are obtained from the position of interface bonding type-E, D and E patterns are obtained from the fracture surface of WC-10Co cemented carbide side when

3.3. Mechanical property test and fracture investigation Fig. 9a illustrates the influence on shear strength under different 5

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Fig. 6. SEM micrographs of five interface bonding types in the weld zone, the EDS results and the SEM micrograph of the weld crack:(a),(c),(e),(g),(i) the SEM micrographs of interface bonding type-A, type-B, type-C, type-D and type-E, respectively; (b), (d), (f), (h) the chemical composition at the interface of interface bonding type-A, type-B, type-C, type-D; (j) the partial enlarged view of the fusion zone of interface bonding type-E; (k) the crack near the welding interface of WC10Co cemented carbide/RM80 steel, which is shown in the red arrow of Fig. 5c. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)

welding currents. Under the welding current 650A (shear stressdisplacement curve shown in Fig. 9b), 800A (shear stressdisplacement curve shown in Fig. 9c) and 950A (shear stressdisplacement curve shown in Fig. 9d), the average shear strength of the welded joints were 700 MPa, 924 MPa and 639 MPa respectively, and are reported in Supplemental information: A. It can be seen that in a certain range, when the welding current increased from 650A to 800A, the joint strength grew and the shear strength increased by 32%. However, after a certain range (I ¼ 950A), the mechanical strength of the joint decreased with the increase of current, and the shear strength decreased by about 31%. Fig. 10g clearly shows that the weld fracture surface presents three different regions at the welding current 650A, which are ductile frac­ ture, brittleness fracture and ductility mixed fracture. The upper part is HAZ (RM80) with deformed dimples and traces of slip (Fig. 10j), indi­ cating this is a ductile fracture. Fig. 11 illustrates that the middle area mainly consists of Fe, Co and W elements, which can prove the white bulk in Fig. 10k is WC particles, and the vicinity of the crack is carbide. It can be considered that the partial fracture occurred in the vicinity of fusion zone, and the fracture mechanism was a mixed one. Traces of slip are not found in Fig. 10l, and this fusion zone basically consists of W and Co elements (Fig. 11d and e), indicating that the fracture of this region occurred on the side of cemented carbide with brittle fracture. It is clear from Fig. 10c and d that the fracture occurred on the steel side at the

welding current 800A, and the EDS analysis result (Fig. 12) shows the fracture surface is substantially made of Fe elements, which can prove that the fracture occurred on the steel side. The fracture surface has deformed dimple feature (Fig. 10m), indicating the mechanism is ductile fracture. When the welding current was 950A, massive WC particles (Fig. 10i) were observed at the middle area of fracture surface after the shear test of sample S3 was carried out, with no dimples being observed. EDS analysis result (Fig. 13) shows that the middle area mainly consists of W elements with a small number of Fe elements, while the sur­ rounding areas mainly consists of Fe elements. It can be concluded that the fracture started from the cracks on both sides and propagate along zone I and zone II, and finally brittle fracture occurred. In summary, if the welding current is too large or too small i.e. when the energy input is either too small or too large, the shear fracture of the welded joints will occur along the loose zone of WC-10Co cemented carbide or near the carbide skeleton of fusion zone (Fig. 10b and f). Under suitable welding current conditions (such as I ¼ 800A), the fracture will occur in the softened zone of HAZ (RM80) (Fig. 10d), rather than near the fusion zone, i.e. the bonding strength of fusion zone exceeds that of the HAZ (RM80).

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Fig. 8. Percentage of the five interface bonding types at the weld interface at the welding current of 650A, 800A and 950A.

Fig. 7. Micro XRD patterns on weld and fracture surface of welded joints, where A pattern is obtained from the position of interface bonding type-D, B and C patterns are obtained from the position of interface bonding type-E, D and E patterns are obtained from the fracture surface of WC-10Co cemented carbide side when welding current is 650A and 950A, while F and G patterns are obtained from the fracture surface of the RM80 steel side.

a brittle layer. However, the shear strength of ultrasonic-assisted brazing is more than 400 MPa, because adding ultrasonic waves can promote the migration of WC particles, increase the thickness of the welded in­ terface’s metallurgical reaction layer, and thereby improving the me­ chanical properties of the welded joints. In addition, it must be pointed out that the brazing is suitable for the relatively complicated welded structural components that generates low manufacturing cost, which is of great significance to engineering manufacturing. In the diffusion welding which adopts different interlayers, the maximum shear strength can reach 620 MPa, which is more than five times the minimum, so using different weld interlayers can also have a large impact on the shear strength. In contrast to fusion welding and brazing, M.N. AvettandFenoel and T. Nagaoka successfully obtained the shear strength of 750 MPa by friction stir welding, because the thermal cycle of the weld only lasted a few seconds, thereby limiting the diffusion of alloying el­ ements and the formation of brittle phase [30]. Similarly, the 924 MPa

4. Discussion 4.1. Joining process and shear strength Table 4 lists the relationships between chemical composition, joining process and the shear strength of cemented carbide/steel joints pub­ lished in recently reported literature. Fig. 14 was plotted based on Table 4 with welding processes as abscissa and the shear strength as ordinate. It can be clearly seen from Fig. 14 that the shear strength of ordinary brazing is lower than 400 MPa, which may due to the fact that the welded interface was exposed to a high temperature environment for a long time, resulting in the formation of massive brittle phase and even

Table 4 Chemical compositions, joining processes, and shear strengths of cemented carbide/steel joints. Process

Cemented carbide

Steel

Maximum shear strength (MPa)

Reference

Brazing Brazing Brazing Brazing Brazing Brazing TIG brazing Vacuum brazing High-frequency induction brazing High-frequency induction brazing Induction brazing Induction brazing ultrasonic-associated brazing ultrasonic-associated brazing Electron beam hybrid welding-brazing Fiber laser welding Diffusion welding Diffusion welding Diffusion welding Diffusion welding Diffusion brazing Rapid diffusion brazing Rotary friction welding Friction stir welding Resistance spot welding

WC–8Co WC-10Co WC-13Co WC–8Co WC–8Co WC–8Co WC-10Co WC – 8Co WC – 8Co WC-15Co WC-15Co WC-20Co WC-15Co WC-15Co WC-30Co WC-20Co WC-15Co WC–8Co TiC-based cermet (TiC, 79 vol%) WC-Co WC-10Co WC-10Co WC-10Co WC-12Co WC-10Co

SAE1045 steel 1Cr13 stainless steel SAE1045 3Cr13 stainless steel AISI 410 carbon tool(0.45%C) AISI 1020 steel tool steel (0.45%C) AISI 4140 35Cr–Mo steel 35CrMo 35CrMo 35CrMo 35Cr–Mo steel 40Cr AISI 1045 90MnCrV8 AISI 410 carbon tool(0.45%C) 90MnCrV8 AISI 4145 40Cr steel AISI 304L S45C steel RM80

203 340 360 154 257 370 289 310 259 350 338 366 408 371 586 980.90(Maximum bend strength) 620 195 380 136 412 293 505 750 924

[36] [29] [37] [14] [38] [33] [10] [26] [39] [40] [15] [17] [27] [40] [9] [21] [41] [42] [43] [44] [28] [32] [11] [30] This work

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Fig. 9. (a) Shear strength of welded samples under different welding currents, (b), (c), and (d) the shear stress-displacement curves of welding currents 650A, 800A and 950A.

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Fig. 10. OM pictures before and after the fracture path at different welding currents: 650A(a and b), 800A(c and d), 950A(e and f); detailed fracture overviews surface observation from SEM imaging of the RM80 steel side: 650A (g), 800A (h), and 950A (i); fracture surface under higher magnification SEM observations: 650A (j, k, and l), 800A(m), 950A(n).

shear strength was obtained by RSW, whose mechanical property is much higher than those reported in the literature. The reasons are as follows: 1) The thermal cycle duration was less than 1000 ms, limiting the diffusion of alloying elements and the formation of bulk brittle phase, and/or 2) under the action of welding pressure, the brittle pha­ se/eutectic phase that will be formed immediately at the welded inter­ face was extruded in the form of slag. The microstructure and fracture modes of the interface will be discussed in more detail in the following

section. 4.2. Formation mechanism of the η phase and interface bonding type According to the liquidus projection of ternary Fe–W–C system (Fig. 15a) and Co–W–C system (Fig. 15b) [45,46], when C content in the system is lower than a certain value, the η phase will be formed. As the surface temperature of the welded joints are only 1600 � C–2000 � C, but

Fig. 11. (a) SEM image of overall fracture at welding current 650A and EDS map of (b) Fe, (c) Ni, (d) Co, (e) W, and (f) C. 9

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Fig. 12. (a) SEM image of overall fracture at welding current 800A and EDS map of (b) Fe, (c) Ni, (d) Co, (e) W, and (f) C.

considering the surface radiation and convection, the partial tempera­ ture inside the welded joint possibly reaches 2800 � C or above. When the temperature of the fusion zone reaches 2785 � C, the edges of the WC particles will begin to passivate [16,47]. The reaction is as follows: (1-1)

WC → W þ C lnCmax ¼ 4:67

heterogeneous nucleation on the surface of WC particles and grow up by “swallowing” the particles [22]. In the actual RSW, due to the rapid solidification of the fusion zone (the whole welding process lasts less than1000 ms, as shown in Figs. 3 and 4), the reaction (1–3) does not have enough time to complete, and the remaining WC particles are surrounded by η-Fe3W3C, forming “spheroidized” WC particles in the fusion zone as shown in Fig. 6e and i, which is also confirmed by Dejian Liu and Liqun Li [45]. Fig. 16 demonstrates the schematic diagram showing the evolution of five interface bonding types, and the schematic view of the fracture path at different welding currents are plotted in the diagram as note in Fig. 17. The experimental results show that the different welding cur­ rents lead to different proportion of five interface bonding types (Fig. 8), which are closely related to the phase transformation mechanism of each zone in the welding process. As shown in Fig. 16b, when the duration at high temperature of fusion zone was relatively short, remelting occurred in parts of areas between the base metals. Therefore, when the welding current was relatively low (I ¼ 650A), the weld interface was occupied by the interface bonding type-A with relatively low bonding strength (700 MPa). This is due to the mild adhesion of cemented carbide and steel and effective welding wasn’t formed; when

15:0 � 103 T

(1–2)

where T is the temperature (in Kelvin temperature), and Cmax is the maximum solid solubility (at%). According to the following equation (1–2) showing the relationship between C atoms and maximum solu­ bility in tungsten [48], the maximum solubility of C in W is only 0.7 at% at 2785 � C. Most of the C atoms will rapidly diffuse into the surrounding liquid phase Co (Fe), causing a “poor C” environment on the surface of the WC particles, in which the following (1–3, 1–4) reactions occur: L þ WC þ W→η–Fe3 W3 C

(1–3)

L þ WC þ W2 C→η–Co3 W3 C

(1–4)

The η phase in the reaction equations (1–3, 1–4) will have

Fig. 13. (a) SEM image of overall fracture at welding current 950A and EDS map of (b) Fe, (c) Ni, (d) Co, (e) W, and (f) C. 10

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Materials Science & Engineering A 776 (2020) 139008

L → γ þ Fe3 C þ WC

5. Conclusions

the welded joint was subjected to shearing force, fracture occurred along the vicinity of the “fusion zone”, whose position is as shown in Fig. 17a. The fracture along the location of the metallurgical reaction layer can be analogized to the model I defined by Gitae Park and Changhee Lee [49] based on interfacial failure (IF). When the temperature reached 2785 � C (Fig. 16c), due to the diffusion of C atoms, the edges of WC particles near the fusion zone dissolved, and the η-Fe3W3C phase was formed on the surface of WC particles (as note in Fig. 6e). With increase of the duration at high temperature, massive Fe atoms “invaded” the matrix of the WC-10Co cemented carbide, causing the matrix to be loose, forming the interface bonding type-C. If a certain zone is exposed to a long duration at high temperature, more complex eutectic reactions (1–5) and (1–6) will occur: ð1380� CÞ

(1–6)

Then the interface bonding type-C were converted to the type-D (Fig. 6g), and the “spheroidized” WC particles in the fusion zone all dissolved and carried out sufficient metallurgical bonding with γ-Fe to form fishbone or herringbone eutectic [50,51]. Under the action of welding pressure or Marangoni convection [52,53] or both, the interface bonding type-C will develop into interface bonding type-E, i.e. forming a relatively large “deformed” fusion zone (as shown in Figs. 6i and 16f)). When the welding current was 800 A and the weld had a metallurgical reaction layer of a suitable thickness, the bonding strength was 924 MPa. The interface bonding type-B, type-C, and type-D accounted for more than 60%, and the fracture path (as shown in Fig. 17b) occurred in the HAZ (RM80) rather than the fusion zone, which was defined as Model II. When the welding current was 950A and the bonding strength was 639 MPa, as the differences in TEC between cemented carbide and steel nearly doubled, under the joint action of excessive thermal cycle and the rapid rise and fall of the temperature, the cracks occurred on the side of cemented carbide (as shown in Fig. 5c), which is the main cause of the deterioration of the welded joints’ mechanical properties. The failure originated from the crack, and the brittle fracture eventually occurred along the Fe–W–C enriched zone of the loose/fusion zone of the cemented carbide. The fracture path is as shown in Fig. 17c with fracture mode III. Therefore, the contribution of five interface bonding types and cracks to the bonding strength of welded joints was determined: type-B and type-C > type-D ≫ type-E > type-A ≫ cracks.

Fig. 14. Average maximum shear strength of different welding methods (BW (Brazing), TIG-BW (TIG brazing), V-BW (Vacuum brazing), HFI-BW (High fre­ quency induction brazing), I-BW (Induction brazing), U-BW (Ultrasonic-asso­ ciated brazing), RFW (Rotary friction welding), DW(Diffusion welding), EB-WB (Electron beam hybrid welding brazing), FSW (Friction stir welding), RSW (Resistance spot welding)).

L → α þ Fe3 W2 þ η–Fe3 W3 C

ð1085� CÞ

In this research, the microstructure evolution, mechanical property and fracture mechanism of WC-10Co/RM80 welded joints under different RSW welding currents were investigated. The main conclusions are summarized as follows: 1. WC-10Co cemented carbide/RM80 steel welded joints were fabri­ cated using RSW process. When subject to annealing treatment at 400 � C/15min in nitrogen atmosphere, the shear strength of the welded joints first increased and then decreased, and the maximum shear strength can reach 924 MPa at pre-welding current of 300A/ 20 ms and welding current 800A/95 ms. 2. With the increase of welding current, the proportion of interface bonding type-A first decreased and then increased, while the

(1–5)

Fig. 15. (a) liquidus projection of the Fe–W–C ternary phase diagram (in Celsius temperature), and (b) liquidus projection of the Co–W–C ternary phase diagram (in Kelvin temperature). 11

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Fig. 16. Schematic diagram showing the evolution of five interface bonding types: (a) pre-weld interface, (b) interface bonding type-A, (c) interface bonding type-B, (d) interface bonding type-D, (e) interface bonding type-C, and (f) interface bonding type-E.

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Fig. 17. Fracture path for the WC-10Co/steel joints at different welding currents:(a) 650A, (b)800A, and (c) 950A

Appendix A. Supplementary data

percentage of interface bonding type-B, type-C, type-D and type-E first increased and then decreased, and cracks were revealed on one side of cemented carbide at the welding current of 950A. The contribution of each interface bonding type and cracks to the bonding strength of welded joints has been established: type-B and type-C > type-D ≫ type-E > type-A ≫ cracks. 3. The shear fracture mode of cemented carbide and steel welded joints can be subdivided into three types. When the welding current was too small or too large, the fracture occurred along the vicinity of the weld. The fracture modes are Model I and Mode III respectively. The former is the fracture without the formation of effective welding, while the latter originated from the crack on both sides of the weld, and then propagated along the loose zone/HAZ (WC–10Co) and finally failure occurred. High quality welded joints will be formed in proper welding current which failed in the HAZ (RM80) with frac­ ture Model II.

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Declaration of competing interest 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. CRediT authorship contribution statement Gang Chen: Supervision, Project administration, Resources, Writing - review & editing. Wei Xue: Conceptualization, Investigation, Writing original draft, Writing - review & editing. Yuzhen Jia: Writing - review & editing, Visualization, Resources. Shucheng Shen: Methodology, Formal analysis, Writing - review & editing. Guoyue Liu: Visualization, Resources. Acknowledgements The authors appreciate the National Natural Science Foundation of China (grant no.51971091) and the emerging industry science and technology project of Hunan Province (grant 2016GK4018) for the financial support to this study.

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