Dissimilar welding of WC–Co cemented carbide to Ni42Fe50.9C0.6Mn3.5Nb3 invar alloy by laser–tungsten inert gas hybrid welding

Dissimilar welding of WC–Co cemented carbide to Ni42Fe50.9C0.6Mn3.5Nb3 invar alloy by laser–tungsten inert gas hybrid welding

Materials and Design 32 (2011) 229–237 Contents lists available at ScienceDirect Materials and Design journal homepage: www.elsevier.com/locate/matd...

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Materials and Design 32 (2011) 229–237

Contents lists available at ScienceDirect

Materials and Design journal homepage: www.elsevier.com/locate/matdes

Dissimilar welding of WC–Co cemented carbide to Ni42Fe50.9C0.6Mn3.5Nb3 invar alloy by laser–tungsten inert gas hybrid welding P.Q. Xu * College of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201615, PR China

a r t i c l e

i n f o

Article history: Received 13 March 2010 Accepted 4 June 2010 Available online 10 June 2010 Keywords: Laser Welding Microstructure

a b s t r a c t Dissimilar welding between cemented carbide and invar alloy was carried out using CO2 laser beam and argon arc as heat sources. g Phase was formed near WC–Co/weld interface and precipitations in the fracture were discovered. In order to disclose the microstructure and mechanical property, firstly, g phase’s morphology and composition at interface were investigated using backscattered electron imaging (BEI); and element diffusion across heat affected zone near WC–Co/weld interface was further studied. Secondly, bend strength values of butt joint with different welding parameters were tested by four-point bend strength experiment. Finally, WC migration mechanism was further discussed and the bend strength was measured. The results showed: (1) microstructures consisted of columnar crystals, cellular crystals, eutectic structure and fir-tree crystal and dendritic crystals. The columnar crystals were surrounded by lots of fir-tree crystals. (2) WC migration was driven by stirring effects of welds, high pressure of molten materials and ionized shielding gas, interface reaction and surface tension. (3) g Phases dispersion did not decrease bend strength of butt joint. And the maximum bend strength was 1493.56 MPa, which was attributed to NbC precipitations featured with super-fine fir-tree. Ó 2010 Elsevier Ltd. All rights reserved.

1. Introduction WC–Co type cemented carbides were used widely in auto industry, mining, die and mould machine, information technology industry, geological prospecting, oil and gas exploration and accurate measurement due to their high hardness (especially in high temperature), high strength with good erosion–resistance and low thermal expansion coefficient [1,2]. The high-end cemented carbide and precision instrument demanded that the deformation kept constant or very small with high strength and good erosion– resistance while it served processing and manufacturing industry. At present, tool made of high speed steel or stainless steel could not meet the requirement of industry. And more and more cemented carbides were needed to substitute them. However, Tungsten and Cobalt were scarce metal, the large-scale use of WC–Co cemented carbides, especial the integral high-end cemented carbide or precision instrument would consume the scarce metals severely. On the other hand, cemented carbides became brittle under severe operational conditions, such as vibration and impact, due to their low fracture resistance and thermal conductivity. Developing cemented carbide weldment and novel no-Cobalt cemented carbide was the two key techniques. Thus, it was necessary to develop

* Tel.: +86 21 677 91474; fax: +86 21 677 91377. E-mail address: [email protected] 0261-3069/$ - see front matter Ó 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2010.06.006

novel welding method to join cemented carbide and steel with perfect formation, well microstructure and high bend strength [3–6]. In order to overcome above problems, dissimilar welding of cemented carbide and invar alloys was put forward using CO2 laser beam and argon arc as heat sources. The incorporation of laser and argon arc not only helped materials improve the absorption of laser power, but also increased the penetration of molten pool [7–9]. Generally, the coupling action of laser and argon arc was valid applying to workpiece consisted of dissimilar materials. The present work aimed at the well formation of penetration of thick cemented carbide and invar steel. This would minimize residual stress, realize full penetration of cemented carbide and contribute to the development of novel cemented carbide product. Interfacial reaction was a key problem when laser was adopted as a heat source. Chen and Li [10] investigated the influence of interfacial reaction on crack initiation and propagation to obtain strong joints. Their experiments indicated that effects of different reaction layer included interfacial layer with insufficient interfacial reaction. The lamella shaped morphology, cellular shaped morphology, club shaped morphology, and thick continuous morphology on crack initiation and propagation were observed. It was found that the joint with lamella-shaped, cellular or serration-shaped, clubshaped interfacial layer had the highest mechanical property. Recently, the welding of cemented carbide to steel had called more and more attention because of the potentially attractive combined properties. Various methods were adopted to join cemented


P.Q. Xu / Materials and Design 32 (2011) 229–237

carbides and steel for drilling and machining tool applications. These methods included braze welding; laser welding; diffusion bonding and mechanical interlock [11–16]. Brazing between cemented carbide and carbon steel using multiple layers of Cu and Ni alloys as insert metal was successfully carried out [12]. The research manifested that as the brazing time was increased; brittle intermediate phases and coarsening of WC grains were produced and deteriorated the joint strength. During this investigation, Co and Cr were emphasized. Co and Cr-rich intermetallic compounds could not be easily formed, because the diffusion of Co from the WC–Co alloy was controlled by the Cr3C2 addition, while higher Cr3C2 content in WC–Co alloy accelerate the formation of carbide during the brazing. Using the similarity principle, VC- and Cr3C2doped WC–NbC–Co cemented carbide was developed [13]. In addition, laser surface melting technique was used to prepare electrometallurgic WC/steel composites [14]. In order to reveal the stress distribution in WC–Co/Cr–Ni multilayer coatings deposited on stainless steel substrate, Sayman et al. [15] carried out thermal stress analysis and finite element model was constructed. By using a thermocouple, the cooling curve of the coating system was measured experimentally and then this curve was considered for numerical analysis. The results showed that the compressive and tensile stresses occurred in WC–Co/Cr–Ni coatings and stainless steel substrate owing to the different thermal expansion coefficients in each material. The effect of carbon on the g phase formation was investigated by Fernandes et al. [16]. The coated powders had revealed a good performance to shape-forming and sintering. However, the presence of nanocrystalline coatings could enhance the formation of a brittle g phase. And the content of g phase was found to be higher for the sputter-coated ones. This result was attributed to the content of binder elements with affinity to carbon, like chromium, enhanced by the difference on binder characteristics. It was concluded that an effective control of the reaction can be achieved by increasing in the binder composition the ratio between the elements without affinity and those with affinity to carbon and/or enlarging the carbon content. Kenichi and Mitsuhiko [17] put forward one efficient estimation method to describe the elasto-plastic-creep characteristics of BAg8 brazing filler metal which was used to weld cemented carbide of K20 (WC–Co) and SKS3 steel. However, porosity and crack were observed during braze welding. In addition, static and dynamic diffusion welding had been adopted to bond cemented carbide and steel using Ni–Cu as interlayer [18]. The fracture strength of welds was markedly improved when Ni interlayer were used. The Ni layer retained at the dissimilar joint interface created a region of weakness. Besides braze welding and diffusion welding, powder metallurgy processing was used to joint cemented carbide and heavy alloys [19]. Wherein, cemented carbide was sinterbonded to nickel–iron tungsten heavy alloy for use in high-temperature tooling applications. Sinterbonding cemented carbide to nickel– iron yielded a consolidated interfaces comprised primarily of complex g phases. In order to minimize the welding stress and decrease the heat affected zone (HAZ) at interface, laser beam welding was a good alternative. Since it was possible to have a control of heat input and obtain limited HAZ [20–22]. Their investigation undertaken with the different laser sources led to full penetration and low cracking welded joints. From the investigation, the addition of laser pulse was help to prevent overheating of cemented carbide and the risk of crack formation. The laser welding of cemented carbide to steel manifested [22] that laser welding allowed the successful joining of steel to cemented carbides without filler. By combining laser welding with a pre-heating and post heat treatment in three-step welding process temperature gradients could be controlled resulting in lower residual stress maximum values and crack-free, non-porous joints. A nitridation of the cemented carbides significantly reduced the amount of

brittle g phase formed in the fusion zone and increased joint bend strength. Among them, the common features were: filler material, groove and heat treatment had been used to relief the welding stress, eliminate the welding crack and inhibit the brittle intermetallic phases by researchers. Nevertheless, the g phases could not be inhibited completely. Because the high temperature in the scope of laser keyhole, laser keyhole shapes were observed experimentally by two highspeed cameras from two perpendicular directions [23]. Their experiment indicated that due to bending of the keyhole, not all the keyhole wall, both on the front and the rear keyhole wall, can be irradiated directly by the laser beam. The laser intensity absorbed on the whole keyhole wall cannot be uniformly distributed even after multiple reflections. Besides the above research, the analysis of WC–Ti, TiC–W laser clad layer was carried out and the results showed that the wear resistance index of a clad layer depended not only on the wear resistance index of the reinforcing phase but also on the bonding strength of the interface between the reinforcement and the matrix [24]. On the basis, in the present work, we addressed microstructure, composition and mechanical properties of welded joint near WC–Co/welds interface. The materials and general experimental procedures were described in Section 2. Section 3 demonstrates microstructure, composition analysis, morphology, Transmission Electron Microscopy (TEM) analysis and desirable elements diffusion near interface. The WC migration mechanism and the effect of welding heat input on the bend strength were further discussed in Section 4; and the conclusions were shown in Section 5.

2. Materials and methods 2.1. Samples description WC–Co cemented carbides (sintering with Fisher Sub-Sieve Sizer (FSSS) 7 lm WC powder and FSSS 2 lm Co binder) were employed as one body material for laser–TIG (Tungsten Inert Gas) hybrid welding. The nominal composition (wt.%) of the cemented carbide was 4.29C, 30Co and 65.71W. The other body material was invar alloy. Its nominal composition (wt.%) was 42Ni, 0.6C, 3.5Mn, 3Nb (balance Fe). The microstructures of cemented carbide and invar alloy were illustrated in Fig. 1. The microstructure of WC–Co (Fig. 1a) consisted of WC strengthening phase and c phase (solid solution W or C) in Co binder. The microstructures of invar alloy (Fig. 1b) included base and precipitated phase. Before welding, the body materials were cut into circular shape (diameter is 48 mm; thickness is 6 mm). The welded joints were produced using a 15 kW CO2 laser (TLF 15000) in combination with TIG welder (MASTERTIG AC/DC 2500). Architecture of laser hybrid welding was illustrated in Fig. 2. Welding parameters were given in Table 1. During laser hybrid welding process, the parameters for laser hybrid welding included laser power, welding speed, gas flow rate, laser spot diameter and defocusing amount. RW in Table 1 was determined by following equation,

RW ¼ g1  U  I=g2  P


where g1 = 0.8; g2 = 0.9, which were the welding heat efficiency of TIG welding and laser welding respectively; P was the laser power; U was the welding voltage; I was the welding current. Pure helium gas was employed as the major shielding gas; and argon gas was used as the minor shielding gas during hybrid welding. The laser spot diameter was 1 mm; and the defocusing amount of laser was 8 mm.


P.Q. Xu / Materials and Design 32 (2011) 229–237 Table 1 Welding parameters during laser–TIG hybrid welding.



1 2 3 4


Fig. 1. Microstructures of body materials: (a) WC–Co and (b) Ni42 Fe50.9C0.6Mn3.5Nb3.

Parameters for TIG welding

Parameters for laser welding

Welding current (A)

Welding voltage (V)

Gas flow (L/min)

Laser power (kW)

Gas flow (L/min)

58 105 149 157

8 12.8 13.8 15.6

Ar, Ar, Ar, Ar,

10 10 10 10

He, He, He, He,

10 10 10 10

32 32 32 32

Welding velocity (m/min)


1 1 1 1

0.04 0.12 0.18 0.22

were then investigated along the position of cross-section and longitudinal direction respectively. Samples for optical and Scanning Electron Microscope (SEM) observations were etched in a solution of mixed acid consisted of 1% HF, 1.5% HCl, 2.5% HNO3, 95% H2O and reagent (15% KOH + 15% K3 [Fe(CN)6] + 70% H2O: Vol.%) at room temperature. Electronic probe microanalysis of sample 2 and 3 was performed using JEOL JXA-8100 Electron Probe Micro analyzer (EPMA); the operating condition was 20 kV and 1 lm diameter. Tungsten, Cobalt, Nickel, Niobium, Iron and Carbon elements were investigated within 500 lm2. Line-scanning analysis was carried out across the WC–Co/welds interface from the weld metal to the unaffected body material. BEI analysis was performed using JEOL JSM-6700 analyzer. Electron Probe Microanalyzer JEOL JXA8100 equipped with an Energy-Dispersive Spectrometer (EDS) was used for morphology observations and composition analysis. A further structural analysis had been carried out using JEM 2010. TEM samples were prepared using standard procedures involving ion milling. The parameters were: voltage: 200 kV; dark current: 96 lA; emission current: 128 lA; the current density: 109.8 pA/cm2; explosive time: 1–8 s and magnification is 20–200k. After polishing treatment, four-point bend strength test was also carried out using dynamic mechanical analysis system. The length of samples was 48 mm curvature radius of the sample was 0.1–0.3 mm, and chamfer angle in head face was 45° c (0.1–0.3 mm).



48mm INVAR





WELDMENT KZ3-C Arc welding automatic controller

Fig. 2. Architecture of laser–TIG hybrid welding process.

2.2. Analyzing position and testing methods For metallographic examination, the samples were cut vertical to welded seam. After grounded, polished and etched; the samples

3. Experimental results 3.1. Macrostructure of welded joints The transverse cross-section of samples 1–4 was illustrated in Fig. 3. While welded with 10 kW laser, welding current I = 58 A, U = 8 V, hot cracks (root crack) were observed (Fig. 3a) near root region which might be induced by recombination action of Ni–Fe chemical system with feature for hot cracking sensitivity and thermal stress caused by contraction distortion during the cooling of welded joint. Phenomenon of insufficient penetration was caused by large physical property difference between WC–Co and invar alloy lack of filler materials. With the increase of heat input of TIG welding, the crack was not observed in sample 2 (Fig. 3b). By contrast, the joints welded with RW = 0.18 (Fig. 3c) presented better formation and shape than the ones welded with RW = 0.04, RW = 0.12 and RW = 0.22. A most wide hyperbolic curve shape profile was observed close to the top surface of welded joint. In the region, the transverse cross-sections showed more wide surface depression and more deep of weld pool caused by dual heat sources. Most of welded joint located in invar alloys side and welding penetration ratio of cemented carbide increased with the increase of heat input ratio of the TIG source to laser power. Contrast to lower heat input welding parameters, humping phenomenon on the surface of the welded joints was observed in


P.Q. Xu / Materials and Design 32 (2011) 229–237







Fig. 3. Macro image of cross-section of (a) sample 1, (b) sample 2, (c) sample 3 and (d) sample 4.

sample 4 (Fig. 3d), which was welded with 10 kW laser, welding current I = 157A, welding voltage U = 14.2 V and RW = 0.22. Close to the top surface of welded joint, a more wide parabolic shape profile was observed. Because the width of welded shape in the middle of the welded joint became larger, the welded joint manifested cone shape. Near the lower welded joint, the cross-section manifested flare shape. The cold cracks, when it happened, always occurred in the cemented carbide; the hot cracks, when it happened, always occurred in the invar alloy. 3.2. Microstructure In the experimental, WC migration phenomenon was discovered far from WC–Co/welds interface, which was shown in Fig. 4a. In addition, the WC migration distance during hybrid welding was much farther than the WC migration distance discovered during TIG welding. WC migration phenomenon happened at the top and root of the cross-section, especially located at the top surface or root of welded joints where full penetration and complete vaporized had happened. With the increase of distance from WC location to top surface, the distance of WC migration became shorter, and in the middle of welded joint, migration phenomenon did not happen. The distance of WC migration in the root could reach 150 lm. And the maximum distance of WC migration across the WC–Co/welds interface could reach 780 lm. The microstructure at WC–Co/weld interface and microstructure of welds were illustrated in Fig. 4b and c. The microstructures near interface were etched by reagent (15% KOH + 15% K3[Fe(CN)6]+70% H2O: Vol.%) at room temperature and consisted of the columnar crystals, eutectic structure, fir-tree crystals. The direction of columnar crystal near interface was vertical to the fusion line which was obvious. The microstructure of welds etched by reagent (1% HF, 1.5% HCl, 2.5% HNO3 and 95% H2O) was made of a lot of cellular structures. The microstructure of welds was different from the microstructures of cemented carbide or invar alloy body

materials. By contrast, the size of microstructures near WC–Co/ welds interface was larger than that of microstructures near invar/weld interface. 3.3. Morphology and composition SEM images of the weld and WC–Co/weld interface were illustrated in Fig. 5, taken near the center of the cross-section of the weldment. The morphology was characterized using backscattered electron imaging and secondary electron imaging respectively. Fig. 5a illustrated SEM images of WC–Co/weld interface. From the morphology, columnar crystals, cellular crystals, eutectic structure, fir-tree crystal and dendritic crystals were observed. The columnar crystals were surrounded by lots of fine fir-tree crystals. The microstructures near cemented carbide interface included bulk tungsten carbide with triangle or tetragonal structure and cobalt matrix. The microstructures near invar alloy interface included cellular crystals, columnar crystals and many super-fine fir-tree crystals. In addition, the dendritic crystals were formed via aggregation of fine crystals. Results on composition analysis of characteristic points near WC–Co/weld interface manifested that average content (wt.%) of Fe in g phase (spectrum 6) was 11.46Fe, 71.18W, 3.27Co, 1.26Co and 12.83C, which was characterized with C depletion. Because average content of Fe and W in M6C and M12C (M stands for Fe, Ni or Co) was nearly 23.2% and 75.30% respectively, g phase belonged to such compound carbides as M6C and M12C. In addition, g phase with W depletion (spectrum 1) was observed simultaneously. The average content of Fe (wt.%) in g phase was 34.45Fe, 29.05W, 24.63Ni, 8.83Co and 3.05C, which was characterized with W depletion. This kind of g phase might be Fe4C2W(Fe–C–W). The average content (wt.%) of fir-tree crystals (spectrum 2) formed on grain boundary in weld consisted of 36.63Fe, 24.85W, 25.43Ni, 10.08Co and 3.01C. Therefore, this fir-tree crystal was NiFe(Co) solution reinforced by WC.


P.Q. Xu / Materials and Design 32 (2011) 229–237



η phase with W depletion

WELD Fir-tree η phase




(b) HAZ




(c) Fir-tree crystals


Fig. 4. Typical WC migration and microstructures near WC–Co/weld interface welded with CO2 laser 10 kW, RW = 0.18, I = 149 A, U = 13.8 V: (a) WC migration (polished), (b) microstructures near WC–Co/weld interface (region ‘B’ in Fig. 3) (etched by reagent (15% KOH + 15% K3[Fe(CN)6] + 70% H2O: Vol.%)) and (c) microstructure of region ‘A’ in Fig. 3 (etched by reagent (1% HF, 1.5% HCl, 2.5% HNO3 and 95% H2O)).

The average contents (wt.%) of cellular crystals (spectrum 3) included 39.04Fe, 16.01W, 31.11Ni, 12.01Co and 1.84C, similar to contents of spectrum 4, which was unbalanced microstructure composed of NiFe solution and WC–Co. The atomic percent for characteristic points of above spectrums was shown in Table 2. Generally speaking, normal phase of W–Co–C system consisted of c + WC, however, when high-carbon phenomenon appeared, carbon content would dissociate from normal W–Co–C system (c + WC) and transited to c + WC + C. On the other hand, when concentration gradient of carbon was large enough, the carbon in c would diffuse from W–Co–C system to weld, followed by diffusion of carbon from WC to c, thus low-carbon phenomenon for W–Co–C system appeared. Subsequently, with the diffusion of Fe or Ni elements, W–X(Fe, Co, or Ni)–C system would appear. From the phase diagram of W–Ni–C, W–Co–C and W–Fe–C system, we

Fig. 5. SEM images of: (a) WC–Co/weld interface, (b) weld, and (c) morphology of intergranular texture near interface.

Table 2 Composition analysis of characterized points (At.%). Element Spectrum Spectrum Spectrum Spectrum Spectrum Spectrum Spectrum

1 2 3 4 5 6 7






38.60 39.84 41.79 47.72 43.90 41.32 2.37

9.89 8.21 5.21 2.46 3.90 40.12 46.49

15.88 15.25 9.14 9.46 6.94 14.18 49.14

26.25 26.31 31.68 28.87 29.04 2.01 2.01

9.37 10.39 12.19 11.48 16.23 2.37 –

knew: scope of normal phase c + WC for W–Ni–C was the most largest zone, followed by scope of c + WC for W–Co–C, and the scope of c + WC for W–Fe–C was the smallest one. Therefore, g phases were easily formed in W–Fe–C system than that in

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W–Co–C even in W–Ni–C system; this was why g phases easily formed with more Fe and less Ni in WC–Co. During laser hybrid welding, in some region near interface, both of WC particles and cobalt matrix melted; in other region, cobalt matrix melted, WC particles did not melt. Therefore, the melt of cemented carbide was attributed to cobalt phases. Consequently, the W contents nearly existed in all the characteristics points, therefore, we knew that WC migration happened across WC–Co/welds interface. 3.4. Element diffusion near interface

200 180

140 120


100 80 60 40

Line-scanning analysis was carried out using BEI method to investigate the element diffusion near interface. The distance was 500 lm; and the intensity was set to 200 cps. The elements included Ni, Fe, C, Co, W and Nb. From the image of line scanning position, the microstructures of cemented carbide and invar alloy were distinct; and the fusion line near WC–Co/weld interface was obvious. The line-scanning result on W, C, Fe, Ni, Nb, Co element diffusion near interface in sample 2 was illustrated in Fig. 6. From it we could see that Nb element existed near interface, but Nb diffusion could not happen; and Fe element diffused obviously. Inharmonic diffusion range of W and Co was observed at the range from 0.1 mm to 0.2 mm. Element W did not decrease until the distance reached 0.2 mm. However, element Co began to decrease at the distance of 0.2 mm. So, while distance was 0.1 mm, the element Co began to melt or evaporate under the effect of laser beam. In contrast with Co atom, Fe, Ni atoms near interface owned strong diffusivity. Content characterized with high Fe would lead to g phase formation, which enhanced opportunity and ability of iron atom to diffuse from weld to W–Co–C system which existed in cemented carbide. And then W–Fe (Ni)–C system (g phase) formed characterized with carbon depletion and Fe atom taking the place of Co atom. Carbon depletion phenomenon of W–Co–C system and high Fe were two vital factors inducing to g phase formation. Meanwhile, evaporating of Co, Fe, Ni led to the inharmonic diffusion of Co and W. The line-scanning result on W, C, Fe, Ni, Nb, Co diffusion near interface in sample 3 was illustrated in Fig. 7. From it we could see that W diffusion happened near interface, and Fe element diffuse obviously. The effects of nickel, iron and carbon on the microstructure and bend strength were as follows, Nickel met the requirement of inhibiting g phase formation with high bend strength. On the other hand, nickel owned the tendency of softening the butt joint

Co Ni W C Fe


Intersity (cps)




Ni Co


C 0.0






Distance (mm) Fig. 7. Line-scanning results for W, C, Fe, Ni, Co diffusion near interface in sample 3.

with lower bend strength; Iron atom diffusion driven by iron content gradient resulted in g phase formation. However, the iron helped to increase bend strength of the butt joint; the carbon decreased the carbon gradient near interface and had the tendency of inhibiting g phase formation. Moreover, carbon strengthened the butt joint. Meanwhile, stress and crack-propagation increased rapidly with increase of carbon. The process of g phase formation was proposed as follows: (a) keyhole wall formed during melting or evaporating of body materials, WC particles moved to weld near interface on the effect of laser beam and argon arc; and then accumulated each other. At solidification stage, element diffusion of Fe, Ni in weld and W, C, Co in WC–Co system happened simultaneously driven by concentration gradient. (b) The effect of laser beam and argon arc was too weak to move the WC particle. WC particles accumulated each other accompanied with interfacial diffusion. At solidification stage, interfacial diffusion continued until g phases formed. Some WC particles melted by laser beam, others did not melt. (c) Aggregation behavior and diffusion behavior underwent simultaneously. WC particles moved to weld near interface on the effect of laser beam, argon arc and interfacial diffusion of atoms driven by concentration gradient until g phases formed. 3.5. TEM results


Intensity (cps)


W C Fe Ni Nb Co



Fe 50

Ni Co Nb C

0 0.0






Distance (mm) Fig. 6. Line-scanning results for W, C, Fe, Ni, Nb, Co diffusion near interface in sample 2.

TEM images of the matrix and precipitation near WC–Co/weld interface were shown in Fig. 8. Wherein, Fig. 8a showed TEM images of the matrix near interface, which was characterized with black and white microstructures with compaction. Fig. 8b manifested that matrix was strengthened by bulk precipitations. The microstructures and its dark-field image of super-fine precipitations were illustrated in Fig. 8c and d. TEM images of the matrix near interface were characterized with black and white microstructures. From the results we could see that the precipitations formed near interface was NbC phase with smaller size. 4. Discussions 4.1. WC migration mechanism Given that the stability of the known binary phases forming in the systems (Fe, Co)–C and W–C was quite different: the tungsten carbide (hexagonal WC) was a very stable compound, while (Fe, Co)3C carbides were thermodynamically unstable and were


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(a) Minipore





Fig. 8. (a) Bright-field TEM image of the matrix observed near interface, (b) bright-field TEM image of matrix far from interface, (c) bright-field TEM image of the precipitation, and (d) its corresponding dark-field image show TEM image of bright-field of matrix and precipitation.

synthesized in a non-equilibrium process [5]. WC migration derived from W–C–Co system on the top surface. Co phases underwent such experience as ‘Vapor–Liquid–Solid’ process. During the process, WC particles in the same region did not undergo ‘Vapor– Liquid–Solid’ experience. The WC migrated far from WC–Co/weld interface driven by forces. In the study, we thought that the stirring action due to the keyhole vapor/liquid interface played a most important role on the WC migration. During laser hybrid welding, intrinsic stress caused by large CTE (coefficient of thermal expansion) gradient between WC and Co was a most significant factor for WC migration. In the experimental, the CTE of WC, NiFe invar alloy and NiFeCMnNb invar alloy was a30–100 = 6.2  106 K1, a30–100 = 4.38  106 K1, and a30–100 = 6.14  106 K1 respectively. However, the CTE of Co was a30–100 = 1.25  105 K1. Therefore, according to the Eshelby model [25,26], there existed lager intrinsic stress from large CTE gradient between WC or invar alloy and Co phase. Different from the body materials, near WC–Co/welds interface, because of the high temperature gradient evoked by heat input, the large intrinsic stress emerged. In addition, during welding processing, the molten cemented carbide at the top surface of WC–Co/welds produced thin jet behind the keyhole resulting in high pressure at the rear end of the molten pool. At the keyhole develops, the power of the source was absorbed at

greater depths not just at the surface [27]. Consequently, on the effect of argon arc, WC migration happened. Therefore, the WC migration was only discovered at the top cross-section of welded joint. Besides the above factors, interface reaction of solid/ liquid interface and liquid/vapor interface was considered also. There existed the forces generated at the liquid/ vapor interface due to the surface tension gradient and also the energy balance at the liquid/ vapor interface and solid/ liquid interfaces. Near WC–Co/weld interface, the Co bond phase and invar alloy underwent Vapor– Liquid–Solid experience and WC particles did not undergo such experience yet. Besides the physical change of WC near interface, interfacial reaction was another factor to lead the WC migration. While different cladding materials were chosen, the crystals near deposited metals/substrate interface manifested different scale, different range, and different morphological characteristics [8,10]. The effect of high energy from laser beam and tungsten arc on the interface made the reaction increase quickly. Wherein, W–Co (Fe, Ni)–C system, rare earth carbide NbC played a key role on the interfacial reaction. The WC, MC1x, M6C, M3C and M2C carbides had been reported to be microconstituents of liquid/solid equilibria [28]. In addition, the WC migration was affected by the hybrid weld pool mixed with strong convection currents resulting from the interaction between the electromagnetic and Marangoni


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forces [29]. The increase in weld pool width was the result of heat input from the arc and the heat transfer aided by strong Marangoni convection. Wider weld pools improved the ability of the hybrid welding process to allow for gaps between large welded sections of material without the necessity of additional bracing. 4.2. Effect of heat input on the bend strength The effect of heat input on the bend strength was carried out using bend strength test and fracture analysis. The results manifested two cases, Case one: to brittle welded joints, the fracture phenomenon would happen; fracture strength could be calculated by following equation,

rbb ¼ 3P  ðL  lÞ=2w  t2


Table 3 The bend strength (rbb, rpb0.2) of different samples. Samples

rbb (MPa)

rpb0.2 (MPa)

1 2 3 4

923.54 1143.91 1483.56 1428.12

– 916.55 1323.54 1137.60


where L stands for the distance between sustain points at the top surface, L = 36 mm; l is the distance between sustain points at the lower surface, l = 10 mm; w stands for width of samples, w = 4.0 ± 0.1 mm. t stands for thickness of samples, t = 3.0 ± 0.1 mm. P is the load; rbb is four-point bend strength. Case two: to ductile welded joints, the whole fracture phenomenon would not happen; fracture strength could be estimated by using rpb0.2 through maximum elastic limits. Bend strength of butt joint varied with different samples welded with different welding parameters, which was shown in Table 3. Sample 1 fractured with lower load because of the root crack and insufficient penetration and manifested brittle fracture. Samples 2–4 were not fractured and manifested ductile property. Fracture morphology of sample 3 was illustrated in Fig. 9. From the fracture image, the crack started from invar alloy, across the weld and terminated at the cemented carbide. Fig. 9a showed that the fracture consisted of WC fracture and welds fracture. The kerfs of WC were observed, and the precipitation around by matrix and boundary could be discovered. And the morphology of precipitation was illustrated in Fig. 9b. We knew that the precipitation NbC was embedded around the matrix. The minor NbC additions refined the WC grains, increased the hardness and bend strength of the base alloys and shortened the thickness of boundary layer [8,30]. The joints produced by laser–TIG hybrid welding had higher bend strength values comparable to those typically observed in laser welded cemented carbide-steel joints [22] or brazed welded joints [31]. g phases were present in the welds, however, they did not form a dense layer near the interface, which had been found to be the main source for a decrease in joint bend strength of cemented carbide–steel joint.

5. Conclusions



Matrix: γ(Fe, Ni)

Precipitation: NbC

Fig. 9. Fracture morphology of: (a) WC–Co/weld interface and (b) precipitation.

1. Dissimilar metals between WC–Co and NiFeCMnNb invar alloy were butt joined successfully by laser–TIG hybrid welding without filler metals. Well-metallurgical joint was obtained while P = 10 kW, I = 149 A, U = 13.8 V, v = 1 m/min, Ar: 10 L/min; He: 32 L/min, free from faults such as porosity (air hole), fine hair crack and slags. 2. The high energy laser beam and argon arc played a key hole on the WC migration across interface. With the increase of heat input, WC migration was more obvious, and the nail head and weld root became larger. WC migration phenomenon happens near top surface and root surface of WC–Co/ weld interface were due to stirring effects of welds, high pressure of molten materials and ionized shielding gas of He and Ar during molten laser keyhole formation, interface reaction and surface tension. In addition, the power driven by the shielding gas and plasma gas dominated the WC migration distance. 3. Two precipitations were observed, one was super-fine fir-tree morphology found in SEM image and TEM image; the other was spherical found in fracture. With the increase of heat input, the bend strength rbb of welded joints increased and then decreased, the maximum value could reach 1493.56 MPa for sample 3. 4. Carbon depletion and diffusion phenomenon were inhibited during the formation of keyhole wall. Under the effect of high energy laser, carbon content dissociated from normal W–Co–C system (c + WC) and transited to c + WC + C. When concentration gradient of carbon increased to some degree, carbon in c diffused from W–Co–C system to weld, followed by diffusion of carbon from WC to c, thus carbon depletion phenomenon in W–Co–C system emerged and led to phase’s transformation of WC ? WnC ? WmMkC.

P.Q. Xu / Materials and Design 32 (2011) 229–237

Acknowledgements The author wish to express thanks to the financially support of National Natural Science Foundation of China (No. 50775135); Natural Science Foundation of Shanghai of China (No. 10ZR1412900); ‘‘Chen Guang” project of Shanghai Municipal Education Commission and Shanghai Education Development Foundation (No. 2008CG62). And many thanks will be given to Professor S. Yao, Y.X. Wu, Dr. F.G. Lu, H.C. Cui from Shanghai Jiao Tong University.

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