International Journal of Minerals, Metallurgy and Materials Volume 25, Number 4, April 2018, Page 439 https://doi.org/10.1007/s12613-018-1589-4

Corresponding author: Mi-qi Wang E-mail: wangmiqihhu@163.com © University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018

Characterization and in-situ formation mechanism of tungsten carbide

reinforced Fe-based alloy coating by plasma cladding

Mi-qi Wang1), Ze-hua Zhou1), Lin-tao Wu1), Ying Ding1), and Ze-hua Wang1)

College of mechanics and materials, Hohai University, Nanjing 211100, China

(Received: 24 September 2017; revised: 24 October 2017; accepted: 27 October 2017)

Abstract: The precursor carbonization method was first applied to prepare W–C compound powder to perform the in-situ synthesis of the WC phase in a Fe-based alloy coating. The in-situ formation mechanism during the cladding process is discussed in detail. The results reveal that fine and obtuse WC particles were successfully generated and distributed in Fe-based alloy coating via Fe/W–C compound powders. The WC particles were either surrounded by or were semi-enclosed in blocky M7C3 carbides. Moreover, net-like structures were confirmed as mixtures of M23C6 and α-Fe; these structures were transformed from M7C3. The coarse herringbone M6C carbides did not only derive from the decomposition of M7C3 but also partly originated from the chemical reaction at the α-Fe/M23C6 interface. During the cladding process, the phase evolution of the precipitated carbides was WC → M7C3 → M23C6 + M6C.

Keywords: precursor carbonization; tungsten carbide (WC); microstructure; in-situ formation mechanism; phase evolution

1. Introduction

The reactive plasma cladding technique has been exten-sively employed in the in-situ synthesis of hard phases in the metal matrix because of its remarkable advantages of rela-tively strong bonding strength, excellent compatibility, and good wettability between hard phases and metal substrate [1]. In previous studies, much effort was devoted to the in-situ synthesis of particles such as TiC [2–3], TiB2 [4–5], ZrB2 [6], and TiN [5,7] on an iron substrate.

Owing to its innate characteristics of extremely high melting point, relatively high hardness, low thermal expan-sion coefficient, and excellent wettability with an iron sub-strate, tungsten carbide (WC) is particularly attractive among various hard phases [8]. However, little research has been conducted on the in-situ formation of WC on an iron substrate using the reactive cladding method. The primary problems that remain unresolved can be understood from two aspects: (1) W is a weak carbide-forming element com-pared to Fe element [9]; hence, finding a method to promote the in-situ reaction between C and W elements in a Fe–W–C system remains an intractable issue. (2) WC is prone to dis-

solve into a molten pool owing to its low free generation enthalpy at high temperatures. Thus, at present we still lack the method to obtain a certain amount of stable WC par-ticles.

The precursor carbonization method is thought to be a very promising method since organics, such as sucrose [10] and asphalt [2–3], are commonly used as both the carbon source and binder of the compound powder. Consequently, each aggregate exhibits a very tight structure comprising fine and homogeneously distributed green materials. In ad-dition, the precursor carbonization method may favor the formation of expected carbide ceramics, which can be as-cribed to the strong adhesive effect of amorphous carbon with other reactive constituents.

Therefore, our study attempted to prepare a W–C com-pound powder using sucrose as the precursor for promoting the in-situ synthesis of WC on an iron substrate during the plasma cladding process. In addition, Fe901 powder was added as a measure against the thermal expansion coeffi-cient gap between the WC and iron substrate which can eas-ily lead to the generation and propagation of cracks. To the best of our knowledge, previous investigations on the reac-

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tion mechanism during reactive plasma cladding are limited. In this study, the in-situ synthesis mechanism of massive carbide generation is discussed in detail.

2. Experimental

Q235 steel was manufactured into a cuboid shape with dimensions of 60 mm × 40 mm × 8 mm as the substrate. A commercial Fe901 alloy powder with a particle size ranging from 15 to 45 μm was chosen as the metal matrix, and its chemical composition is listed in Table 1. 99.5wt% tungsten powder with an average particle size of 2.6 μm was selected as the green material. Sucrose was used as the precursor of the carbonization process.

Table 1. Chemical composition of Fe901 powder wt%

Cr Si B Mo Fe

13 1.2 1.6 0.8 Bal.

The carbonization temperature was determined by the thermal behavior of the W–C compound powder, as shown in Fig. 1. In this study, since the mass loss tended to be sta-ble at 873 K, this temperature was selected as the carboniza-tion temperature. Moreover, the calculated carbonization ra-tio was approximately 23.1%.

The W–C compound powder was prepared via a novel process. First, the mixtures of tungsten and sucrose were heated at 623 K for 30 min under an argon atmosphere to facilitate the evaporation of substances and impurities with low molecular weight. Subsequently, the mixed powder was carbonized at 873 K for 2 h. Finally, the compound powder was cooled down in the furnace.

Fig. 1. TG-DSC curves of W–C compound powder.

The cladding materials consisting of the Fe901 and W–C compound powder were pre-coated on the substrate. Then, the layer was dried at 373 K for 2 h. The plasma cladding procedure was performed under argon protection on

LHD-300-type equipment. The optimum parameters were 1.2 for the atomic ratio of C to W, 20wt% for W content, 130 A for cladding current, 100 mm/min for cladding speed, 0.8 L/m3 for plasma gas flow, and 1.2 L/m3 for protective gas flow.

Thermal analysis was conducted using a differential scanning calorimeter (NETZSCH STA 409 PC) from room temperature to 1073 K under an argon atmosphere and a heating rate of 10 K/min. The phase composition of coatings was identified using X-ray diffraction (XRD, BRUKER D8) with Cu Kα radiation. The cross-section microstructure was observed via a scanning electron microscope (SEM, Hita-chi-S3400N) equipped with X-ray energy dispersion (EDS). The microhardness of the cross section was tested under a load of 0.98 N in 15 s using a Vickers hardness tester (HXD-1000TC).

3. Results and discussion

3.1. Microstructure and microhardness of composite coating

Fig. 2 shows the SEM micrograph of the W–C com-pound powder prepared using a precursor carbonization method. It can be seen that the reactive powder has a par-ticle size of up to 10 μm and exhibits a polygonal shape. Moreover, the compound powder exhibits an exceedingly tight structure because of the adhesive effect of amorphous carbon. It was assumed that the tight structure provided by the precursor carbonization method could increase the in-situ reaction areas between W and C elements.

Fig. 2. W–C compound powders with a tight structure pre-pared using precursor carbonation method.

The phase composition of the composite coating is shown in Fig. 3. The relatively strong diffraction peaks cor-responding to WC can be easily identified; this suggests that the expected WC phases were successfully generated when Fe/W–C compound powder was used. Simultaneously, no peaks were consistent with W, which substantiates that the

M.Q. Wang et al., Characterization and in-situ formation mechanism of tungsten carbide reinforced … 441

in-situ reaction was sufficient. Various unavoidable carbides, such as M23C6, were also observed from the XRD pattern. This will be further discussed as follows.

The cross-section morphology of the composite coating is shown in Fig. 4. As shown in Fig. 4(a), a deep, dark pla-nar growth region with a thickness of 10 μm appeared at the coating/substrate interface and provided a good metallurgic-al bonding that was obtained at the interface [11]. As the distance from the coating/substrate interface increased, a 50-μm-thick cellular growth region was observed. The evo-lution of the solidified structure can be attributed to the change of G/R [7], where G is the temperature gradient at the solid-liquid interface, and R is the solidification speed.

Fig. 3. XRD pattern of composite coating.

Fig. 4. Microstructure of composite coating: (a) interface region; (b) lower region; (c) upper region; (d) magnified reinforcements.

Massive reinforcements were densely distributed in both the lower and upper regions of the composite coating in Figs. 4(b) and 4(c). The reinforcement microstructure is further verified in Fig. 4(d). Fig. 4(d) shows a magnified image of typically dispersed reinforcements, and the corresponding EDS results are listed in Table 2. It should be noted that the composite coating was free of cracks and pores from the low and high magnification photographs. Various fine and white bright particles (marked as A) were either surrounded by or were semi-enclosed in a blocky phase (marked as B). Moreover, herringbone-like (marked as C) and net-like (marked as D) precipitations were also observed. In addition, the black matrix (marked as E) was attached by the above-mentioned precipitates.

The white bright phase was determined as WC since the

chemical composition was 51.02at% C and 49.98at% W, in which the atomic ratio of C to W was close to 1. It was in-teresting to discover that the WC particles exhibited an obtuse shape with no sharp angles, rather than a hexagon-al or triangular shape, as has been reported previously in Refs. [12–13]. The primary elements in the blocky phase were Fe, W, and C, in addition to a trace amount of Cr. This blocky phase was possibly M7C3 because of the presence of 30at% carbon content. It is well known that Cr7C3 carbides are an interstitial compound [14], in which Cr atoms are easily substituted by other atoms such as Fe and W.

The herringbone-like phase was enriched in Fe (45.94wt%) and W (44.95wt%) elements and was con-firmed as M6C-type carbides. The occurrence of the her-ringbone growth of M6C-type carbides showed a tendency

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of grain coarsening, which may deteriorate the composite coating hardness. Based on EDS analysis, the chemical com-position of a net-like structure was identified as 65.55Fe– 8.04Cr–6.51W–19.90C (at%), which is a typical eutectic composition of α-Fe and M23C6. It was inferred that α-Fe was at the core, whereas the net-like shell was mainly M23C6.

Table 2. EDS analysis results of composite coating obtained by reactive plasma cladding wt%

Region C Cr Fe W

A 6.37 — — 93.63

51.02 — — 49.98

B 4.37 5.81 26.54 63.59

28.09 8.62 36.70 26.59

C 4.07 5.03 45.94 44.95

22.54 6.44 54.74 16.27

D 4.33 7.58 66.38 21.71

19.90 8.04 65.55 6.51 E — 4.84 81.30 13.86

The black region, primarily composed of the Fe element and a small amount of Cr and W, was verified as α-Fe solid solution. The dissolution of Cr and Win the α-Fe matrix strengthened the metal matrix and was beneficial to the wear resistance of composite coating. In terms of distinguishing each phase, the microhardness of five morphologies was measured and is presented in Table 3. Among them, it is clear that the coarse M6C-type carbides expressed the lowest microhardness.

Table 3. Phase compositions and microhardness of five mor-phologies

Region Morphology Possible phase Microhard-ness, HV

A White bright region WC 1725 B Blocky region M7C3 1522

C Herringbone-like

region M6C 1228

D Net-like region M23C6 and α-Fe 1351 E Black matrix α-Fe 1393

3.2. In-situ formation mechanism and phase evolution

In our study, the following reactions in the Fe–W–Cr–C system may emerge by integrating the experimental results to the previous discussion, during the formation of multiple products:

L α-Fe→ (1)

W + C WC→ (2)

7 3L + C M C→ (3)

7 3 6L + M C M C + α-Fe→ (4)

7 3 23 6L + M C M C + α-Fe→ (5)

23 6 6M C + L M C + α-Fe→ (6)

First, Fe901 powder melted into a metal liquid (L) due to its low melting point. A subsequent reaction between W and C occurred according to Eq. (2). A previous study has pointed out that W2C tends to precipitate from the molten pool owing to its more negative Gibbs free energy of forma-tion compared with that of WC at high temperature [15]. However, no single phase is identified as W2C in our study. The excessive moles of carbon may only account for the appearance of WC.

As the primary phase, WC is prone to precipitation ow-ing to its extremely high melting temperature. With the nuc-leation and growth of WC particles, the region around them was rich in Fe and Cr elements. Furthermore, a certain amount of C and W atoms, released by WC particles, dis-solved into the molten pool and gathered with Fe and Cr near the surface of WC to form M7C3. It can be speculated that WC could provide the nucleation sites for M7C3-type carbides. Therefore, M7C3 carbides were attached to the edges of WC particles and led to WC particles either being surrounded by or being semi-enclosed in M7C3 hard phases, as shown in Fig. 4 (d).

With the cooling of molten pool, metastable M7C3 car-bides served as another primary phase. Owing to their low decomposition temperature of approximately 873 K [16], stable carbides, including M23C6 and M6C, were in-situ transformed from M7C3 by consuming free atoms from the metal liquid. The nucleation of various M6C herringbone carbides also occurred at the α-Fe/M23C6 interface. In this case, the source of W atoms was provided by M23C6 because of its higher W content compared to that in M6C, as shown in Table 2. Moreover, for a higher C content of M6C com-pared to that of M23C6, the C element in M6C was not only derived from M23C6 but also from the dissolution and diffu-sion of C atoms at the α-Fe/M23C6 interface [17]. Thus, the phase evolution during the reactive plasma cladding process was determined as WC → M7C3 → M23C6 + M6C.

3.3. Analysis of obtuse WC particle formation

The relation between the equilibrium melting point and the curvature of the solid phase interface during the growth and solidification process of the WC particles [18] is given below:

m sr

2T V RT

H

(7)

where r mT T T is the difference to the equilibrium melting point, Tm is the equilibrium melting temperature, T is the actual melting temperature, Vs is the molar volume of solid particles, R is the average surface curvature of solid particles, σ

M.Q. Wang et al., Characterization and in-situ formation mechanism of tungsten carbide reinforced … 443

is the interfacial tension, and ΔH (negative value) is the molar change in the enthalpies of the solid/liquid transition.

For the WC solid particles, the curvature of the sharp an-gle on the solid phase surface is a positive value. Thus, ΔTr is a negative value based on Eq. (7). In other words, the melting point of the sharp angle will reduce. The larger the curvature, the lower is the equilibrium melting temperature. Therefore, sharp angles with large curvatures will be the first to dissolve in the metal liquid, which will lead to a smooth marginal. However, since the cooling rate of plasma cladding can reach up to 105°C/s, fine and obtuse WC par-ticles can be retained well.

4. Conclusions

(1) W–C compound powders prepared using the precur-sor carbonization method showed a very tight structure, which can increase the bonding strength of the W/C inter-face and promote the in-situ reaction to generate WC.

(2) Fe-based coatings reinforced with WC and other complex carbides were successfully synthesized from the Fe/W–C powders. The composite coating exhibited good metallurgical bonding with the substrate.

(3) Fine WC particles with polygonal shapes were dis-persed in α-Fe matrix. The appearance of obtuse WC may be attributed to the lower melting point of sharp angles at the marginal.

(4) During the cooling process of the molten pool, in-situ synthesized WC phases experienced a transformation into blocky M7C3 that generated around the edges of WC. Fur-thermore, the decomposition of M7C3 at high temperatures led to the occurrence of eutectic net-like structures com-posed of M23C6 and α-Fe. Moreover, coarse M6C in her-ringbone shape originated from the decomposition of M7C3 and the dissolution reaction at the α-Fe/M23C6 interface.

Acknowledgements

This work was financially supported by the National Natural Science Foundation of China (No. 51379070).

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