Vacuum 178 (2020) 109467

Available online 19 May 20200042-207X/© 2020 Elsevier Ltd. All rights reserved.

The investigation on friction and wear properties of cast iron matrix surface compact vanadium carbide layer

Lisheng Zhong a,b,*, Shaoxiong Zhang a,b, Xin Wang a,b, Haiqiang Bai a,b, Jianlei Zhu a,b, Zhengxin Lu a,b,c, Yunhua Xu a,b

a School of Material Science and Engineering, Xi’an University of Technology, Xi’an, 710048, PR China b International Research Center for Composite and Intelligent Manufacturing Technology, Xi’an University of Technology, Xi’an, 710048, PR China c Xi ’an Zhitong Automation Technology Ltd. Co., Xi’an, 710048, PR China

A R T I C L E I N F O

Keywords: Vanadium carbide In-situ solid-phase diffusion Microstructure Friction Abrasive wear

A B S T R A C T

To enhance the wear resistance of cast iron, a cast iron matrix compact vanadium carbide layer was fabricated by two step in-situ solid-phase diffusion processes. The vanadium carbide layer consist of V2C dense layer, V8C7 dense layer, and V8C7/α-Fe gradient transition layer with 7.6 μm, 35.7 μm, and 86.3 μm in thickness, respec-tively. The friction and wear properties of the compact vanadium carbide layer and the cast iron substrate were evaluated using the pin-on-disc abrasive wear tests and ball-on-flat dry sliding wear tests. Abrasive wear results show that the wear resistance of the compact vanadium carbide layer is 11.67 times higher than that of matrix. The main abrasive wear mechanism is fracture and peel off of vanadium carbide particle. Ball-on-flat dry sliding wear results show that the friction coefficient of the layer increased first and then decreased with the temperature changing from 25 �C to 800 �C. The main friction wear mechanism is spalling at 25 �C, forming mechanical mixture layer at 400 �C, and delamination wear at 800 �C.

1. Introduction

Composites coated with transition metal carbide reinforcements have high melting point, high bulk modulus, and high hardness and have received considerable attention for potential application in metallurgy, mining, machinery, printing and other field [1,2]. Because the transition metal carbide reinforcements, e.g. TiC [3], WC [4], ZrC [5], TaC [6] and NbC [7], have excellent oxidation and corrosion resistance. Therefore, these transition metal carbides are usually used as the surface coating of the metallic materials. Iron and its alloys are widely used as matrix [8]. The reasons are that iron and its alloys have low cost, good machin-ability and weld ability, whist those are low strength and poor wear resistance.

Numerous preparation technologies have been reported in preparing carbide coating reinforced iron and its alloys, including thermo-reactive diffusion treatment [9–11], magnetron sputtering [12], laser cladding [13], and several other methods [14]. In particular, the composites coated by vanadium carbide reinforcements show good wear resistance [10–13]. The hardness and wear resistance of vanadium carbide rein-forced with a Fe-based composite coating are improved due to the

presence of in situ vanadium carbide particles fabricated by laser clad-ding [13]. The thickness of vanadium carbide layers range from 2.3 mm to 23.2 mm, which prepared on steel by thermal reactive diffusion. The hardness of the coating is approximately 2487 HV; these coatings exhibit superior wear properties after pin-on-disk abrasive wear test [10].

However, there is less investigation on the two important aspects which are discussed as following. First, vanadium carbides exhibit different phase types, chemical compositions, and crystalline structures, i.e., V2C, V4C3, V6C5, V8C7, and VC. The hardness, strength, and elastic modulus of vanadium carbides are gradually improved when the C/V ratio is increased [15–17]. Furthermore, the C/V ratio of V8C7 is 0.875. And V8C7 is more thermodynamically stable than VC [18]. Hence, V8C7 is more suitable as reinforcements for carbide reinforced metal matrix composite. Second, because of the limitation of preparation process, the volume fraction of vanadium carbide in the composite is limited. Abrasive wear resistance of the low volume fraction of vanadium car-bide coating material on tool steel was poor against SiO2 particles, which prepared by PTA-welding method [14]. Conversely, the high volume fraction of vanadium carbide compound layer on CI can be fabricated by in situ reaction and the detail process has been discussed in previous

* Corresponding author. School of Material Science and Engineering, Xi’an University of Technology, Xi’an, 710048, PR China. E-mail address: zhonglisheng@xaut.edu.cn (L. Zhong).

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Vacuum

journal homepage: http://www.elsevier.com/locate/vacuum

https://doi.org/10.1016/j.vacuum.2020.109467 Received 1 April 2020; Received in revised form 11 May 2020; Accepted 12 May 2020

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works [19]. Although several studies obtained new insights on preparing composites coated with vanadium carbide reinforcements, the friction and wear properties have received little attention, especially for the compact vanadium carbide (V8C7) layer. Therefore, the wear mecha-nisms of V8C7 and wear properties of vanadium carbide coating with high volume fraction should be in-depth and detailed research; in addition, the effect of experimental temperature on the change of wear mechanisms is also described.

In this study, a compact vanadium carbide layer on cast iron (CI) was prepared through a two-step in situ solid-phase diffusion (ISSD) process. In order to simulate the actual application conditions, the abrasive wear and friction properties of the layer were studied by pin-on-disc abrasive wear tests and ball-on-flat dry sliding wear tests under normal atmo-spheric conditions, respectively. The relative wear resistance, friction coefficients and worn morphology were observed and analyzed; and the relationship between the characteristics of V8C7 carbide, experimental temperature and wear mechanisms were investigated and discussed in detail.

2. Experimental procedures

2.1. Specimen preparation

Cast Iron (CI) and vanadium (99.99% purity, 1 mm in thickness) were used as the raw materials in this experiment. The chemical com-ponents of CI were analyzed using a direct reading spectrometer and were 3.45C, 0.56 Si, 0.268 Mn, 0.024 S, and 0.224 P (wt. %). The compact vanadium carbide layer on CI was prepared through a two-step ISSD process. The production procedure of the sample was similar to that introduced in a previous study [19]. The schematic diagram of the specimen preparation process is shown as Fig. 1. First, the CI was placed on top of the vanadium plate contacted with the polished surface, and then the stacked couple was placed into a vacuum furnace under vac-uum degree of 10� 3 MPa. The samples were heated from room tem-perature to 1110 �C for 5 min; then cooled to room temperature. After that, the samples were reheated to 1000 �C for 2 h and cooled to room temperature. The heating rate and the cooling rate of heat treatment are 6 �C/min and 5 �C/min, respectively. For contrasting, the contrast sample were produced with the a tube furnace under argon atmosphere with the same heat treatment process, and the argon flow rate is about 5 mL/min. Vacuum environment and inert argon gas were used to isolate samples from air and protect samples from oxidize and nitride.

2.2. Characterization

The cross-section and surface of the specimens were polished and etched with Nital. The phases were tested by X-ray diffractometer (XRD, Japan), and the XRD patterns were recorded in the 2θ range of 30�–90�at the scan rate of 5�/min. The microstructures of the specimens were examined by scanning electron microscopy (SEM, ZEISS, Germany).

Electron backscattered diffraction (EBSD) analysis is performed by Oxfard-EBSD system; the imaging working distance is 10 mm, the ac-celeration voltage is 20 kV, and the sample is tilted at 70�. The field- emission transmission electron microscopy (TEM, JEM-3010) was car-ried out to characterize the microstructure and selected area electron diffraction (SAED) patterns. The micro-hardness of the specimens was measured using a micro-hardness tester (HDX-1000, Cewei, China). The static load and dwell time were 0.1 kg and 10 s, respectively.

2.3. Pin-on-disc abrasive wear tests and ball-on-flat dry sliding wear tests

The specimens were cut to Ø6 � 10 mm for the pin-on-disc abrasive wear tests. The relative wear resistance (β) was utilized to estimate the pin-on-disc abrasive wear resistance of the compact vanadium carbide layer. This value can be obtained by

β¼ΔVr=ΔVc (1)

where ΔVr is the wear volume of the CI, and ΔVc is the wear volume of the compact vanadium carbide layer on CI. The worn surface measure-ments were taken by scanning electron microscope (SEM, JSM-5800, JEOL, Japan).

Ball-on-flat dry sliding wear tests were carried out using tribometer (HT-1000, Zhongke, China). The Si3N4 ball with 5 mm diameter was employed to slide against the samples (Φ20 mm � 10 mm). Each wear test keeps 30 min (180 m), a load of 10 N, and a speed of 0.1 m/s. Each sample was tested three times at a relative humidity of 60%–70% and at room temperature (25 �C), 400 �C, and 800 �C.

The weight loss of the flat cylinder was measured and converted to volume loss values as mentioned above. The wear tracks were observed by SEM. The XRD data of the worn surface were recorded by X-ray diffractometer (PW 1730, Philips, Netherlands) with Cu Kα radiation at 40 kV and 40 mA.

3. Results and discussion

3.1. Microstructure and micro-hardness of vanadium carbide layer on CI

Fig. 2 shows the microstructure of the cast iron (CI) matrix surface vanadium carbide layer fabricated by vacuum furnace and argon at-mosphere furnace, respectively. A “layered þ gradient” transition zone is formed between metal vanadium and CI by both processes, which are consistent with previously reported [19]. All the organization from the vanadium plate to the vicinity of CI is: V2C dense layer, V8C7 dense layer, V8C7/α-Fe gradient transition layer. Therefore, these two preparation conditions that guarantee a non-oxidizing environment have basically similar microstructures. For the narrow transition layer in Fig. 2b, it is caused by the temperature variation of different furnace. For the oxide or nitride contaminated by residual gas, it was not found in the formed layer of samples which produced by two kinds of furnace. But the former research results confirmed that the low vacuum degree (0.05 MPa)

Fig. 1. The schematic diagram of the specimen preparation process.

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would result in the increased oxygen content in formed layer [20]. Therefore, when the vacuum degree is higher than 10� 3 MPa, the question of residual gas contamination can be avoided.

In this paper, the mainly focus is the structure and properties of the samples prepared by vacuum furnace. As shown in Fig. 2a, close to va-nadium is a very thin V2C dense layer with approximately 7.6 μm in thickness. Near the CI is a V8C7/α-Fe gradient transition layer with approximately 86.3 μm in thickness. Between the V2C dense layer and V8C7/α-Fe gradient transition layer exists a V8C7 dense layer of approximately 35.7 μm in thickness. Different microstructures entail different mechanical properties. The “layer þ gradient” transition zone can also be analyzed by testing the micro-hardness distribution. From the V2C dense layer to the CI, the micro-hardness decreased from 2034 HV0.1 to 235 HV0.1, thereby showing a gradient distribution. Given the dense short-range V2C, the highest hardness occurred at the V2C dense layer. The average value of the V2C dense layer is up to 1963 HV0.1, which is eight times that of CI (235 HV0.1) and four times than that of metal vanadium (505 HV0.1). For the same reason, the average hardness of the V8C7 dense layer is up to 1707 HV0.1. For the V8C7/α-Fe gradient transition layer, the average hardness is 735 HV, which is approximately three times that of CI. The hardness distribution corresponds to the microstructure.

The residual metal vanadium can be ground off through sandpaper before the detail phase analysis. Fig. 3 shows the XRD results of the vanadium carbide layer. The characteristic peaks are identified as V8C7 and α-Fe phases. V8C7 (JCPDS 35-0786) is a substoichiometric carbide where the vacancies of carbon is ordered. The lattice parameter of V8C7 is a ¼ 0.8334 nm [21]. However, the V2C phase was not detected, which may be because the thinner V2C layer was ground away during the preparation of the surface sample.

Fig. 4 shows the EBSD results of the surface of the vanadium carbide layer on CI. The dense V8C7 grains are distributed on the surface. The α-Fe phases are distributed around the V8C7 grains. No obvious V2C phase was found in the whole layer. This result is consistent with the XRD results (Fig. 3). The grain size of the V8C7 grains was analyzed with Tango software. The average grain size of the V8C7 particles was 0.96 μm, and the area fraction of V8C7 carbide is higher than 95.1%.

The TEM was adopted to characterize the microstructure and the crystal structure of the particles in vanadium carbide layer. The results are shown as Fig. 5, in which the Diff.1 indicates that the dark particle has BCC (Body-Centered Cubic) structure; further the interplanar spacing and angles are in accordance with the crystal structure of α-Fe. Meantime, the Diff.2 indicates that the light particle has a FCC (Face- Centered Cubic) and the interplanar spacing and angles are in accor-dance with the crystal structure of V8C7. Combined with the results of XRD, EBSD and the former research results [19], it can be concluded that the dark particle and light particle is α-Fe and V8C7, respectively. Fig. 5 (a) reveals that the interface between the V8C7 phase and α-Fe phase exhibits a good micro-interface. And according with the phase distri-bution map of EBSD (Fig. 4), α-Fe phase is distributed along the grain boundary of the V8C7 particles.

3.2. Pin-on-disc abrasive wear

Fig. 6 shows the wear loss, relative wear resistance and wear mor-phologies of the vanadium carbide layer on CI under continual abrasive wear. From the wear loss curve, the wear loss trend of the layer on CI is in accordance with the hardness distribution. The relative wear resis-tance of the V8C7 dense layer on CI is 11.67. Of note, the wear loss of V8C7/α-Fe showed a gradient increase. The wear pattern of the V8C7 dense layer shows little damage. Little peel off and fracture of the car-bide particles were observed. Conversely, deep and wide ploughings that show serious wear were observed on the wear pattern of CI after abrasive wear test. The relationship among the wear rate, Q; volume of a groove, Vg; and length of the groove, L, can be expressed as [22].

Q¼ ηVg=L (3)

where η is fraction of the material displaced from the groove that is effectively removed as wear debris (η ¼ 0, pure microploughing; η ¼ 1, pure microcutting).

The abrasive particle dragged across the surface of CI produces a groove, which is shown in Fig. 7a. The volume of a groove is given by [22,23].

Vg¼Lhc2 tan φ (4)

hc¼

�2

πtan2φ⋅FN

H

�1 =2

(5)

Fig. 2. Cross-section microstructure and micro-hardness distribution of the vanadium carbide layer on CI under vacuum (a) and argon atmosphere (b).

Fig. 3. XRD results of the surface of the vanadium carbide layer.

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where hc is the indentation depth, ϕ is the abrasive particle character-istic angle, FN is the normal load, and H is the hardness.

Combining Eqs. (3)–(5) results in

Q¼ 2ηFN�

πHtan2φ (6)

The relation implies that the wear rate decreases with the increase in the hardness of the material. The average micro-hardness of the V8C7 dense layer is 1707 HV0.1 which is approximately seven times higher

Fig. 4. EBSD analysis results on the surface of the vanadium carbide layer: (a) quality map and (b) phase distribution map.

Fig. 5. TEM analysis of the surface of the vanadium carbide layer on CI: (a) bright field TEM image; (b) and (c) SAED patterns of areas Diff. 1 and Diff. 2 indicated in panel (a).

Fig. 6. (a) Wear morphologies of CI (b)wear morphologies of the V8C7 (c)Wear loss and relative wear resistance of the vanadium carbide layer on CI under continual abrasive wear.

Fig. 7. Schematic of the abrasive wear behavior of CI (a) and the vanadium carbide layer (b).

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than that of CI matrix. The wear pattern of the CI shows deep and wide ploughings (Fig. 6a), with almost pure microcutting, η ¼ 1 (Fig. 7a).

For the V8C7 dense layer, the particle size of the V8C7 is larger than the indentation depth because of the increase hardness of the compact V8C7 (Fig. 7b). The V8C7 dense layer had high hardness, but the V8C7 particulates were not able to resist the reciprocal action of aluminium oxide, which was observed in the surfaces (Fig. 6b). Such outcome in-dicates neither pure microploughing nor pure microcutting, 0 < η < 1. Therefore, the V8C7 dense layer was moderately worn, the wear rate decreased, and the relative wear resistance increased to 11.67.

3.3. Ball-on-flat dry sliding wear

3.3.1. Wear rate and friction coefficient Fig. 8 shows the wear rate and friction coefficient of the vanadium

carbide layer and CI at room temperature (25 �C), 400 �C, and 800 �C. The friction coefficient of the tested material increases with the sliding time of up to 2–3 min due to pre-grinding. During this time, no signifi-cant change occurs on the friction coefficient. The friction coefficients of the V8C7 dense layer on CI at 25 �C, 400 �C, and 800 �C were 0.53, 0.64, and 0.20, respectively, whereas those of CI were 0.65, 0.90, and 0.41, respectively. Clearly, the coefficient of the V8C7 dense layer and CI increased first and then decreased with increase in temperature.

The wear rate results show that the wear rate of the V8C7 dense layer increased with the increase in temperature. The wear rate of the V8C7 dense layer increased rapidly when the temperature was higher than 400 �C. For CI, the trend of the wear rate was opposite to that of the V8C7 dense layer. The increase of the wear rate of CI slowed down with the increase in temperature. The relative wear resistance β of the V8C7 dense layer was 3.9, 7.6, and 2.7 at 25 �C, 400 �C and 800 �C, respectively.

3.3.2. Wear characteristics Fig. 9 shows the XRD results of the worn surface of the compact layer

under 25 �C, 400 �C and 800 �C. At room temperature, no oxidation occurred on the surface. The main phases of the surface of the vanadium carbide layer were V8C7, α-Fe, and V2C. After sliding at 400 �C for 30 min, obvious V3O4 characteristic peaks were observed on the surface. When the temperature reached 800 �C, the V3O4 peaks strengthened and increased.

Fig. 10 and Fig. 11 show the wear morphology and schematic of the sliding dry wear behavior of the V8C7 dense layer and CI substrate at 25 �C, 400 �C, and 800 �C, respectively. At room temperature, the surface of the V8C7 dense layer after wear was relatively flat due to the good co-ordination between the V8C7 particles and α-Fe. Only non-obvious furrow and pits arose on the surface of the V8C7 dense layer (Fig. 10a). Pits were caused by the spalling of V8C7 particles under cycle compressive stress and shear stress (Fig. 11a). Under compressive stress, the α-Fe distributed among the V8C7 particles underwent plastic defor-mation. When the Si3N4 ball started to slide, different deformation de-grees between α-Fe and V8C7 produced the stress concentration at the interface. Under the cycle compressive stress and shear stress, micro-cracks appeared at the interface. The spalling of V8C7 particles occurred due to the growth of multiple microcracks or interconnect. Therefore, the main wear characteristic of the V8C7 dense layer is spalling at 25 �C. However, the main wear characteristics of CI are abrasive wear and adhesive wear (Fig. 10b).

At 400 �C, the α-Fe bonds to the V8C7 particles and the shear strength of the V8C7 dense layer are reduced due to high temperature [24]. For one thing, the spalling characteristic became serious due to oxidation, fatigue and work hardening. The oxidation products here are V3O4, Fe2O3, and Fe3O4 [25]. For another, the mechanical mixture layer

Fig. 8. Friction coefficient and wear loss of the vanadium carbide layer and CI substrate at room temperature (25 �C), 400 �C and 800 �C.

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(MML), which is composed of abrasive debris, the broken carbide par-ticulates, and oxidation product, is formed due to their re-inserts into the material (Fig. 11b). Hence, the wear loss of the V8C7 dense layer is relatively gentle from 25 �C to 400 �C (Fig. 8c). Thus, the main wear characteristic of the V8C7 dense layer is MML due to oxidation and re-inserts of the oxidation product, abrasive debris, and the broken carbide particulates at 400 �C. For CI, abrasive wear and adhesive wear became serious due to the reduction of hardness on the surface of the V8C7 dense layer.

At 800 �C, obvious delamination wear arose on the wear morphology of the V8C7 dense layer due to severe oxidation at high temperature. The oxidation products here are V3O4 and FeO [25]. For one thing, cracks were generated in the interface of the oxidation product layer and V8C7 dense layer during the sliding. Then, the cracks expanded and linked up with other cracks until reaching the worn surface, finally resulting in delamination wear (Fig. 11c). For another, after the oxidation product layer peeled off, new V8C7 particles were exposed in the environment and oxidized. Therefore, the wear loss of the V8C7 dense layer increased sharply from 400 �C to 800 �C (Fig. 8c). Thus, the main wear charac-teristic of the V8C7 dense layer is the delamination wear at 800 �C. For CI, large scale spalling pits can be observed on the wear morphology. This outcome is due to the extensive spalling of FeO oxidation layer at high temperature.

4. Conclusions

A cast iron matrix surface compact vanadium carbide layer was successfully fabricated through a two-step in-situ solid-diffusion reac-tion process. The abrasive wear and friction properties of the layer were investigated in detail. The following conclusions could be drawn:

(1) A “layered þ gradient” transition zone is formed between metal vanadium and cast iron, which consists of a V2C dense layer, V8C7 dense layer, and V8C7/α-Fe gradient transition layer with a thickness of 7.6, 35.7 and 86.3 μm, respectively. After removing the residual metal vanadium, the surface was dominated by V8C7 and α-Fe, and the area fraction of V8C7 carbide is higher than 95.1%.

(2) Pin-on-disc abrasive wear test results show that the relative wear resistance of the compact layer is 11.67. The main abrasive wear mechanism is the carbide particle fracture and peel off. Ball-on- flat dry sliding wear test results show that the relative wear resistance of the layer is 3.9, 7.6, and 2.7 times higher than that of CI at 25 �C, 400 �C, 800 �C, respectively.

(3) The friction coefficient of the V8C7 layer first increased and then decreased with the temperature change from 25 �C to 800 �C. The main friction wear mechanism is spalling at 25 �C, MML forma-tion at 400 �C, and delamination wear at 800 �C.

Fig. 9. XRD results of the worn surface of the vanadium carbide layer on CI under 25 �C, 400 �C and 800 �C.

Fig. 10. Wear morphology of the vanadium carbide layer and CI at 25 �C (a and b), 400 �C (c and d) and 800 �C (e and f).

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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.

Acknowledgements

The work was supported by the National Natural Science Foundation of China (grant no. 51974242) and the Key-point Research and Inven-tion Program of Shaanxi Province (grant no. 2017ZDXM-GY-043). The authors also gratefully acknowledge the funding support of the Inno-vation Capability Support Program of Shaanxi (grant no. 2018GHJD-17 and 2019-TD019).

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