In situ fabricated metal-carbide with core–shell structure for high impact-toughness iron-matrix composite

Abstract

A series of metal-carbide (Ta–TaC, Nb–NbC and W–WC) with core–shell structure for iron-matrix composites are fabricated by in situ solid-phase diffusion. Results show that the formation of metal-carbide with a rod-shaped core–shell structure, in which the metal-rod surface was covered with a carbide shell layer, in the iron- matrix after in situ solid-phase diffusion. The TaC, NbC, and WC shell layers are in situ synthesised by the diffusion of carbon atoms from the iron-matrix onto the surface of the Ta, Nb, and W rods, respectively. Metallurgical integration occurs between metal-carbide and iron-matrix. The metal-carbide-reinforced iron-matrix composites show excellent impact resistance, and the shell-layer hardness is extremely high.

Keywords

Carbide core–shell structure Iron-matrix Composites Diffusion Impact resistance Hardness

Introduction

Carbide-reinforced steel-iron-matrix composites (CSMCs) combine the excellent toughness of steel-iron-matrix and the high hardness, strength, modulus, high-temperature durability, and low thermal expansion of carbide. Thus, CSMCs are widely used in many abrasion-resistance applications. Some common carbides used in such composites include tungsten carbide (WC) [1 –4], titanium carbide [5 –7], vanadium carbide [8 –10], chromium carbide [11], tantalum carbide (TaC) [12,13], and niobium carbide (NbC) [14]. Among these CSMCs, TiC particle reinforced steel-iron matrix composites is currently the most widely used [5 –7,15 –17]. For example, steel-based cemented carbides with the registered trademarks Ferro-TiC, Alloy-TiC and Ferro-Titanit are widely used as wear-resistant materials and high-temperature structural materials, and their properties are significantly better than existing tool steels [15 –17]. Recently, the development trend of the preparation method for CSMCs is in situ synthesis. Compared with the ex situ method, in situ synthesis has many advantages [18 –20], including a generally clean and compatible interface between carbide and steel-iron matrix, thermodynamically stable and low degradation in high-temperature applications of in situ synthesised carbide, and strong interfacial bonding between carbide and metal. Accordingly, many studies have focused on in situ synthesis methods, such as casting [4,21], plasma-transferred arc [10,22], selective laser melting or laser cladding [1,7,8,11], hot-pressed sintering [12], spark plasma sintering [3,4], vacuum sintering [6,13], and arc melting [2,14].

Generally, CSMCs show excellent wear resistance. However, their toughness, yield strength, and maximum compressive strength may decrease to some extent with increased carbide fraction [8,12]. For example, Peng et al. [12] fabricated TaC particle-reinforced iron-matrix composites with 20–60 wt-% TaC content through high-energy ball milling followed by subsequent hot pressing. With increased TaC content from 20 to 60 wt-%, the yield strength and maximum compressive strength of the composites initially improve markedly and then decrease, whereas the strain decreases with increased TaC content, especially for 60 wt-%-TaC/Fe composite. One of the reasons is that the stress concentration area increases with increased carbide volume fraction, and the other is that cracks easily form due to the large difference in physical properties (e.g. wettability, elastic modulus, and thermal expansion coefficient) between carbide and iron- matrix. Moreover, the applications of carbide are restricted because carbide in the steel-iron matrix is very difficult to distribute uniformly due to the large difference in density between carbide and steel-iron matrix. For example, the density of TaC (14.3 g·cm⁻³) is more than two times that of steel [12,13]. Although the density of NbC (8.47 g·cm⁻³) is similar to that of steel, preparation results show that partial segregation and uneven distribution are inevitable [14].

To overcome these shortcomings, metal-carbide with a rod-shaped core–shell structure were designed in the present study. Dense carbides were distributed around corresponding metal rods. Iron-matrix composites reinforced with the core–shell structure metal-carbide were prepared by in situ solid-phase diffusion method. In this structure, carbides were ‘segregated’ around a metal rod, which improved the hardness and wear resistance of the composites due to higher the higher volume fraction of carbide. The metal ‘core’ with excellent toughness further positively contributed to improving the toughness of the composites.

Material and method

Cast iron (94.044Fe–3.210C–0.045S–0.077P–0.014Cu–0.240Cr–1.050Mn–1.320Si) with dimensions of 12 mm × 12 mm × 60 mm was used as the carbon source for in situ synthesis. Rods of pure metal, including Ta (diameter, 2 mm; 99.99% purity), Nb (diameter, 1.5 mm; 99.99% purity), and W (diameter, 0.8 mm; 99.99% purity), were used as the raw materials. The core–shell structure metal-carbide-reinforced iron-matrix composites were produced by in situ solid-phase diffusion method. First, a 2 × 2 array of Φ2 mm × 60 mm, Φ1.5 mm × 60 mm, Φ0.8 mm × 60 mm holes was drilled on a 12 mm × 12 mm surface of the cast-iron substrate. Then, the metal rods and cast iron were ultrasonically cleaned in alcohol for 20 min. Second, metal rods were inserted into the cast iron to form a precursor. Here, the metal rod was tightly fitted to the hole in the cast iron to promote atomic-diffusion transmission. Third, the precursor was placed in a horizontal tube furnace (X1200, Kejing, China) with an argon gas flow rate of 10 ml min⁻¹ at 1150°C for 80 min (Ta rod), 1135° for 60 min (Nb rod), and 1100° for 60 min (W rod), respectively. The specimen was cooled room temperature naturally. Herein, the in situ solid-phase diffusion reaction temperature for the preparation of the core–shell structure Ta–TaC-reinforced iron-matrix composites was selected to be 1150°C for the following reasons. (i) The temperature was determined based on the differential scanning calorimetry results reported in our previous studies [23]. The ternary eutectic temperature of the Ta–Fe–C system was 1165°C. (ii) Based on the Fe–C binary phase diagram, the binary eutectic temperature of the Fe–C system was 1148°C. Theoretically, at 1150°C, the cast iron was in a semi-solid state and only partially melted, which also contributed to the diffusion of carbon atoms at the interface. (iii) At 1150°C, a rapid diffusion of carbon atoms was anticipated, resulting in high productivity. Previous studies have shown that the ternary eutectic temperature of the W–Fe–C system is 1084°C [24,25]. At 1100°C, cast iron (T_m = 1148°C) and W rod (3410°C) remained solid. However, a W–Fe–C ternary system molten pool formed at the interface between the cast iron and W rods, which promoted atom diffusion at the interface and improved interface bonding strength. Thus, the in situ solid-phase diffusion reaction temperature for the preparation of the core–shell structure W–WC-reinforced iron-matrix composite was selected to be 1100°C. For the core–shell structure Nb–NbC-reinforced iron-matrix composite, the tight distribution of NbC particles around the Nb rods were controlled by an in situ solid-phase reaction temperature of 1130°C based on our previous study [26].

To examine phase composition and microstructure, the specimens were wire-cut into dimensions of 12 mm ×12 mm ×15 mm. The 12 mm ×12 mm surface of the specimens was polished with diamond paste and ultrasonically cleaned in acetone. The phase composition of the surface of specimens was determined by X-ray diffraction (XRD; Philips, Netherlands) with monochromatic Cu K_α radiation at 40 kV and 40 mA in the 2θ range of 20°−90°. XRD data were analysed using the International Centre Diffraction Data. The polished surface of the samples was etched with 4% nital for microstructure analysis, which was performed using scanning electron microscopy (SEM) system (JSM-6700F, JEOL, Japan) equipped with an energy-dispersive X-ray spectrometer (EDS). The micro-hardness (H) of the iron-matrix composite was measured with an HDX-1000 digital micro-hardness tester (G&R Technology Inc, USA), which consisted of a square-based pyramidal diamond indenter with a 136° angle between two opposite faces. The static load was 50 g and the dwell time of loading was 15 s. The Charpy impact test was applied to measure the impact toughness of the iron-matrix composite at room temperature using a JB-300B impact tester (Jinan Puye Electrical and Mechanical Technology Co., Ltd., China). The size of the impact toughness specimen was 10 mm × 10 mm × 55 mm. The orientation of the metal rods in the composite during the impact test was perpendicular to the direction of impact. The compression performance of the iron-matrix composites was measured with a computer servo control universal material testing machine (HT-2402) with a compression rate of 0.3 mm min⁻¹. The cylinder with a structure unit of metal-carbide core–shell rod (Ф3 mm × 6 mm) was cut from the prepared composite. The orientation of the metal rods in the composite during the impact test was parallel to the direction of impact. The average value of impact toughness and compression performance was obtained from three different measurements. The fracture morphology was characterised by SEM.

Results and discussion

According to the ternary phase diagram [4,27 –29], no liquid phase existed in the Ta–Fe–C, Nb–Fe–C, and W–Fe–C ternary systems at 1150°C, 1135°C, and 1100°C, respectively, which were selected as the in situ solid-phase diffusion temperatures. However, the temperature provided sufficient energy for carbon atoms to undergo interstitial diffusion from the iron-matrix into the metal. Thus, solid-state phase changes occurred and carbide formed by nucleation and growth [30]. Furthermore, a compact carbide shell layer formed on the metal-rod surface (i.e the core–shell structure metal-carbide). The XRD pattern of the core–shell structure Ta–TaC-, Nb–NbC-, and W–WC-reinforced iron-matrix composites are shown in Figure 1. The XRD analysis revealed that the TaC, NbC, and WC phases were present as new phases within the Ta-cast iron, Nb-cast iron, and W-cast iron precursors, respectively. These phases markedly enhance the hardness of composites in other studies [1 –4,12 –14]. From the perspective of thermodynamics, TaC, NbC, and WC with Gibbs free energies of −6.28, −40.80, and −34.98 kJ·mol⁻¹ were the most stable phase of Ta–Fe–C, Nb–Fe–C, and W-Fe–C ternary systems at 1150°C, 1135°C, and 1100°C, respectively [23,24,26]. From the perspective of dynamics, carbon atoms with small atomic radius can easily cross the diffusion energy barrier at high temperature [31]. The enrichment of carbon on the metal-rod surface provided a prerequisite for the formation of carbides by solid-state phase changes.

Figure 2 shows the macro- and micro-structure of the core–shell structure Ta–TaC-, Nb–NbC-, and W–WC-reinforced iron-matrix composites. For the Ta–Fe–C system, the outer surface of Ta rod was covered with a dense layer of carbide particles, which was denoted as the core–shell rod, after heat treatment at 1150°C for 80 min. The diameter of the core–shell rod was extended to about 2.1 mm compared with the 2 mm in diameter of original Ta rod. According to the EDS results obtained from the carbide particles present in the shell layer, the ratio of Ta to carbon was 50.48:49.52 or near 1:1, in accordance with the XRD result. For the Nb–Fe–C system, after heat treatment at 1135°C for 60 min, a dense layer of carbide particles surrounded the Nb rods. The diameter of the core–shell rods was extended by about 1.6 mm compared with the 1.5 mm diameter of the original Nb rods. According to the EDS results, the dense carbide particles content Nb and C. The ratio of Nb to carbon is 50.93:49.07, near 1:1, which means it is NbC particles combine the XRD results. As well, for the W–Fe–C system, the diameter of the core–shell structure W–WC extend to about 1.0 mm compared with the original 0.8 mm diameter original W rods after heat treatment at 1100°C for 60 min. With respect to the formation of WC, a previous study [32] reported that the carbon concentration was required to be approximately 3.0 and 6.13 wt-% for the synthesis of W₂C and WC, respectively. Yuan et al. [33] prepared WC reinforced iron-based composites by in situ metallurgical reaction using W, C and Fe-30wt-%Ni alloy powders as raw materials. Studies have shown that the phase composition were WC and M₆C, while W₂C and M₁₂C were not found. In addition, Cai et al. [34] fabricated a WC–Fe cemented carbide layer on the surface of the iron substrate by in situ reaction. The phase composition was WC and α-Fe phases. Notably, the W₂C, M₆C, and M₁₂C phases were not detected in the W–WC-reinforced iron-based composites, which may be due to the following two reasons. (i) In situ solid-phase diffusion had a high carbon potential due to the high concentration of interstitial carbon atoms in the cast iron, which provided a strong thermodynamic driving force for the diffusion. (ii) The phase content in the composite was relatively low, and the XRD identification of the type of phase is difficult at <5% phase content.

The TaC, NbC, and WC layers can be in situ synthesised through the diffusion of carbon atoms from the iron matrix onto the surface of the corresponding Ta, Nb, and W rods by heat treatment at a lower temperature near the ternary eutectic temperature. On one hand, metallurgical integration led to the formation of defects such as porosity, slag inclusion, and microcracks, between the core–shell structure metal-carbide and iron-matrix. On the other hand, differences existed in the particle size and morphology of carbide with different metal rods. For example, TaC particles in the core–shell structure Ta–TaC were cubic with no more than 1 μm particle size, similar to previous results [12]. For core–shell structure Nb–NbC, NbC were nearly spherical or irregular block-shaped particles with 1–3 μm particle size, also similar to previous results [35]. For core–shell structure W–WC, WC had particles that were large and were irregularly shaped, slightly different from previously reported WC morphologies [33]. The discrepancies may be due to the differences in insulation time. A shorter insulation time prevents complete WC growth [33]. Furthermore, all carbide shell layers were not composed entirely of pure carbides and a small amount of α-Fe existed, which may have helped improve the toughness of the iron-matrix composites.

As aftermentioned above, the main purpose of designing and preparing the core–shell structure metal-carbide-reinforced iron-matrix composites was to obtain high strength and toughness. In the composite, the strength was enhanced by dense carbide, which was distributed around the corresponding metal rods, whereas the toughness was improved by the metal ‘core’ with excellent toughness. Table 1 shows the mechanical properties of CSMCs prepared by different methods. Numerous studies have focused on hardness and only a few pay attention to the toughness and strength of CSMCs. The core–shell structure metal-carbide-reinforced iron-matrix composites show excellent impact toughness and high strength. In the present study, the toughness of the core–shell structure Ta–TaC-reinforced iron-matrix composites was evaluated by impact tests and determined to be approximately 100.6 J·cm⁻² at room temperature. To understand the toughening mechanism, the fracture morphology was characterised and results are shown in Figure 3. The fracture morphology of the metal Ta ‘core’ is displayed in Figure 3(a). The Ta ‘core’ showed a clear necking, indicating its excellent toughness. Moreover, the interface between the TaC shell layer and the Ta ‘core’ exhibited a strong bond, indicating that the fracture was along the TaC/TaC (or TaC/α-Fe) interface (as shown in Figure 3(b)). Figure 3(c) displays the fracture morphology of the TaC–Fe shell layer as a typical cleavage brittle fracture attributed to the high volume fraction of the hard brittle TaC phase. Figure 3(d) presents the fracture morphology of the cast iron matrix, indicating its excellent toughness. Furthermore, the compression performance of the core–shell structure Nb–NbC-reinforced iron-matrix composite was evaluated by compression tests, as shown in Figure 4. Figure 4(a) shows the compressive stress–strain curves of the of the core–shell structure Nb–NbC-reinforced iron-matrix composite. The maximum compressive strength and strain are 2238 MPa and 11.8%, respectively. Figure 4(b,c) presents the fracture morphology of the NbC–Fe shell layer. Secondary cracks, which can induce the initial cracks to deflect and blunt, can be observed (Figure 4(b)). More NbC particles were pulled out from the fracture surface, indicating that crack initiation and propagation occurred along with the NbC/Fe interface, particularly intergranular fracture. The cast iron matrix further exhibited a typical dimple fracture morphology.

Table 1

The mechanical properties of NbC reinforced steel-iron-matrix composites prepared by different methods.

Ref.	Method	Matrix	Enhanced type	Hardness	Yield strength (MPa)	Fracture (MPa·m^1/2)
[36]	Liquid phase sintering	M48 high speed tool steel	Particle	14.91 ± 0.14 GPa	–	1.56 ± 0.3
[37]	Powder metallurgy	Iron	Particle	139 HV_0.5	235	–
[38]	Centrifugal casting	High chromium white cast iron	Gradient	489 HV	–	20.8 ± 0.54
[39]	Vacuum sintering	AISI 440C stainless steel	Particle	86.2 HRA	1987.9 (TRS)	–
[40]	Vacuum sintering	Vanadis 4 tool steel	Particle	815 HV	2093.3 (TRS)	–
[41]	Vacuum sintering	Stainless steel	Particle	13.6 GPa	–	7.3
[42]	Vacuum sintering	Hadfield steel	Particle	87 HRA	–	17
This work	In situ solid-phase diffusion method	Iron	Particle Core-shell	1988 HV_0.05 (shell)	1981	5.89 (shell)

Overall, the high strength and toughness of the core–shell structure metal-carbide reinforced iron-matrix composite can be attributed to the synergistic effects of the high-volume-fraction carbide ‘shell’ layer and the high-toughness metal ‘core.’ First, the residual metal ‘core’ with excellent toughness played an important role in improving the toughness. Metal Ta, Nb, and W are known to have excellent plasticity and toughness [43], so strong plastic deformation occurred and most energy was absorbed by the metal ‘core’ under an applied load. These phenomena played important roles in improving the toughness of the iron-matrix composites. Second, the modified iron matrix was more conducive to improving the toughness of the composite. The in situ solid-phase diffusion process consumed and reduced the carbon content in the iron matrix. With decreased carbon content, the toughness of iron matrix increased. Third, a small amount of α-Fe in the carbide ‘shell’ positively contributed to improving the toughness of the composite. The high volume fraction of the bundled hard phase carbides in the ‘shell’ layer significantly improved the strength of the composite.

Conclusion

The impact toughness of CSMCs decreased with an increased fraction of carbide to some extent. A series of core–shell structure metal-carbide-reinforced iron-matrix composites was prepared by in situ solid-phase diffusion method. First, the structure of the core–shell structure metal-carbides included metal-rod surfaces covered by a dense carbide layer. This layer was formed through carbon atom diffusion from the iron-matrix onto the metal-rod surface. Given the dense and small size of carbides in the ‘shell’ layer, a small amount of α-Fe existed. Second, the core–shell structure metal-carbide can be fabricated at a low temperature near the ternary eutectic temperature, i.e. 1150°C, 1135°C, and 1100°C for the Fe–Ta–C, Fe–Nb–C, and Fe–W–C systems, respectively. Third, metallurgical integration led to the lack of defects such as porosity, slag inclusion, and microcracks between the core–shell structure metal-carbide and iron matrix. Fourth and last, the iron-matrix composites reinforced by the novel core–shell structure showed an excellent impact toughness at room temperature of about 100 J^.cm⁻², and the hardness of the TaC shell layer was 1832 HV_0.05 on average. The residual metal ‘core’ with excellent toughness and the modified iron matrix were conducive to improving the toughness of the composite. The small amount of α-Fe in the shell layer also made a positive contribution.

Footnotes

Disclosure statement

No potential conflict of interest was reported by the authors.

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