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Effect of boronizing temperature and time on microstructure and abrasion wear resistance of Cr12Mn2V2 high chromium cast iron
Surface & Coatings Technology 202 (2008) 5882–5886
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Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Effect of boronizing temperature and time on microstructure and abrasion wear resistance of Cr12Mn2V2 high chromium cast iron Chunmin Li, Baoluo Shen ⁎, Guijiang Li, Chao Yang College of Materials Science and Engineering, Sichuan University, Chengdu 610064, China
A R T I C L E
I N F O
Article history: Received 16 January 2008 Accepted in revised form 12 June 2008 Available online 27 June 2008 Keywords: High chromium cast iron Boriding Microstructure Abrasion wear resistance
A B S T R A C T In this study, Cr12Mn2V2 high chromium cast iron (HCCI) was boronized at 900 °C and 950 °C for 2, 4, 6, and 8 h, respectively. The borided samples were investigated by X-ray diffraction (XRD), scanning electron microscopy (SEM), microhardness tester and ring-on-block wear tester in terms of the phase composition, microstructure and worn surface morphology, microhardness, fracture toughness and abrasive wear resistance. The boride layer thickness ranges from 8 to 33 µm. XRD studies show the boride layer formed at 950 °C/8 h consists of the phases FeB, Fe2B and CrB, while for the layer formed at 900 °C/8 h, consists of mainly the Fe2B phase. Abrasive wear test results show that the relative wear resistance of the borided HCCIs increases with increasing surface microhardness. © 2008 Elsevier B.V. All rights reserved.
1. Introduction High chromium cast iron is widely used for wear components in the mining, mineral and oil industry due to its high hardness (because of the presence of hard (Fe,Cr)7C3 carbide) [1–2]. However, the cylinder liner that is made of HCCIs in mud pumps is very susceptible to damages in oil drilling machines. Improving the lifetime of such component is an important task. In recent decades, research activities have focused on surface treatments such as boriding, chromium coating, ion implantation and so on [3]. Boriding is thought to be an important thermochemical treatment to enhance the surface hardness and wear resistance of ferrous and non-ferrous alloy components [4]. Solid-state boriding has some important advantages in terms of easy handling, the flexibility with respect to the composition of the powder, minimal equipment and low cost [5–7]. The aim of this paper is to carry out solid-state boriding of Cr12Mn2V2 cast iron and investigate the resulting microstructure and properties of the borided Cr12Mn2V2 high chromium cast iron using X-ray diffraction (XRD), scanning electron microscopy (SEM), microhardness tester and ring-on-block wear tester. 2. Methods and procedures 2.1. Substrate materials Cr12Mn2V2 hypoeutectic HCCI samples were used for this study, with the following composition (wt.%):2.76%C, 12.62%Cr, 0.90%Si,
⁎ Corresponding author. Tel.: +86 13688456619; fax: +86 02885402231. E-mail address: [email protected] (B. Shen). 0257-8972/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2008.06.170
1.98%Mn, 1.6%V, 0.072%P, 0.036%S, and balance Fe. The HCCI was cast from 1500 °C as 80-mm-diameter balls into a chilled mould. Then the balls were cut into a size of 10 mm × 10 mm × 10 mm samples for the boriding treatments. Before the boriding treatment, all the samples were ground with 600 grid emery papers to get the final surface finish and then ultrasonically cleaned in acetone. 2.2. Treatments Boriding treatment was carried out in a sealed stainless steel container located in an electrical resistance furnace. The container was filled with the granular boriding medium that consisted of the boriding source B4C, the activator KBF4 and the diluent SiC. The test samples to be borided were placed in contact with the boriding medium. Then, boriding was performed at the constant atmosphere at 900 °C and 950 °C for 2, 4, 6, and 8 h, respectively, followed by furnace cooling. 2.3. Characterization The microstructure of the polished and etched cross-section of the specimens was observed by scanning electron microscopy (SEM, JSM5910LV). The thickness of boride layer was also measured using SEM. The Type Rigaku D/Max-rA X-ray diffraction apparatus was employed to examine the phase composition of the boride layer, with the radiation Cu-Kα, the tube voltage 42 kV and the tube current 110 mA. The microhardness profile of the boride layer was assessed by Type SHIMADZU microhardness tester with a load of 25 g. The surface microhardness and the fracture toughness of the boride layer were measured by Type MVC-1000A microhardness tester with the applied load 500 g and 300 g, respectively. The fracture toughness calculation was obtained by the equation Kc = 0.028(E/Hv)1/2 · P/c3/2[7], where E is
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Fig. 3. Cross-section view of the HCCI borided at 950 °C for 8 h. Fig. 1. Schematic representation of a Vickers indention in a coated material.
the Young's modulus of the boride layer (approximately 290 GPa), Hv is the layer hardness, P is the applied load, and c is the indentation half crack length as shown in Fig. 1. The abrasive wear test was carried out on a Type M200 ring-onblock wear tester. A schematic view of the tester is shown in Fig. 2. The test was performed at a constant speed of 200 rpm (equal to a liner velocity of 52.3 cm/s) under a load of 30 N. For the test duration, 10 min was selected .The counterface ring is a 50-mm-diameter 45#steel ring pasted with the 120#-alumina (Al2O3) cloth. Mass loss (M) of the sample was measured by TG328A photoelectric balance with accuracy of 0.1 mg. The wear rate was defined as mass loss per unit sliding distance (mg/km) and the relative wear resistance was defined as the ratio of wear rate of the uncoated cast iron (but held at 1000 °C for 0.5 h followed by air cooling) to that of borided cast iron. The worn surface morphology of the borided sample was also observed by SEM. 3. Results and discussion 3.1. Microstructure Fig. 3 shows the SEM image of the etched cross-section of the HCCI borided at 950 °C for 8 h. As can be seen, a compact and relatively smooth layer is formed on the surface and no obvious transition zone
Fig. 2. Schematic diagram of the abrasion tester.
is found. The morphology of the boride layers depends fundamentally on the chemical composition of the substrate. It is reported that sawtoothed layers are formed on the low-alloy steels while the layers are flat in the case of high-alloy steels [7,8]. As substrates, the Cr12Mn2V2 HCCI has a higher content of alloying elements, such as Cr, Mn and V, among which, Cr and V help with the formation of the FeB phase but inhibit the growth of saw-toothed layer. In the treating process, these atoms tend to concentrate at the tips of the boride columns, reducing considerably the active boron flux in this zone. Consequently, the columnarity was decreased at the interfaces [6]. In addition, the interface between the boride layer and the substrate is not obviously distinguished. This is probably due to the combination of borides and carbides. Fig. 4 shows the boride layer thickness of the HCCIs as a function of time. It can be observed that the thickness is parabolically increasing with time and its thickness value is in the range of 8–33 µm. As it is well known, thickness of the boride layer is closely related with the process temperature, treatment period and chemical composition of substrate [9]. Due to the fact that the process temperature plays a much more important role in the layer thickness increase than the treatment duration [7], the boride layers formed at 950 °C show larger thickness than that at 900 °C as shown in Fig. 4. On the other hand, it was found that the boride layer thickness of the Cr12Mn2V2 sample is much
Fig. 4. Boride layer thickness of the borided HCCIs.
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Fig. 7. Surface microhardness (HV0.5) of the HCCIs treated at different borided conditions. Fig. 5. XRD patterns of the borided HCCIs.
smaller than that of the other low and medium carbon plain carbon steels including low alloyed steels [6,9]. This phenomenon can be explained as follows: (1) Higher carbon content in the HCCI: since the carbon is not soluble in iron borides, it can be concentrated beneath the boride layer while diffusing into the substrate. So the boride diffusion was prohibited and the formation of the boride layer was hindered [10]. (2) The large amount of Cr and V in the HCCI substrates can reduce the active boron diffusivity by entering the iron boride lattice, as a results, the boride layer thickness tend to decrease [7]. 3.2. XRD analysis Fig. 5 reveals the XRD patterns of borided samples. As shown in Fig. 5, it can be observed that the diffraction peaks of the Fe2B produced on the surface of the 900 °C/8 h boronizing sample and 950 °C/ 8 h boronizing sample are high and sharp, however, the diffraction peaks of the FeB and CrB formed at 900 °C/8 h become weak or disappeared in comparison to that formed at 950 °C/8 h. This suggests that lower process temperature is favorable to form a single phase Fe2B for the HCCI. The reason for this is most probably attributed to the lower boron potential in the boriding process. Though there is 1.6% V
in the HCCI, no vanadium borides (e.g. VB) were observed from the results of all the XRD patterns. This is probably because vanadium borides are too little to be detected. 3.3. Microhardness and fracture toughness Fig. 6 shows the microhardness profile measured on the crosssection of the HCCI borided at 950 °C/8 h. It can be seen that the average value of microhardness of the borides reaches 1566 HV0.025, which is much greater than that of the substrate due to the present FeB, CrB and Fe2B phases in the boride layer. Moreover, it can also be observed that the microhardness fluctuated within a narrow range .This can be attributed to the presence of structural defects (e.g. porosity, cracks, etc.) and the different type of the borides [4]. Fig. 7 shows the surface microhardness(HV0.5) of the HCCIs treated at different borided conditions. As can be seen, the surface microhardness of the borided samples varies between 719 HV0.5 and 1650 HV0.5. It increases with boriding time extending and temperatures increasing. The FeB phase was found to be much harder than Fe2B phase [8,11], and Ugur Sen also showed [12] that the Vickers microhardness of the FeB phase varied between 1920 and 2140 HV0.1, the inner Fe2B phase ranged from 1160 to 1920 HV0.1 in the borides formed on the ductile iron substrate, so the boride layers formed at 950 °C show an obvious increase in hardness compared with those at 900 °C due to the FeB phase formation. Moreover, as a result of the large load (500 g) applied, borides formed on the HCCIs show an obvious phenomenon of indentation size effect (ISE, which usually involves an apparent decrease in the microhardness with the increasing applied test load) [13].Therefore, both the hardness of the single phase Fe2B and the indentation size effect are the dominant reason for the lower surface microhardness of the borides formed at 900 °C. Indentation fracture toughness tests are performed for the samples borided at 900 °C and 950 °C for 6 h and 8 h. The fracture
Table 1 Fracture toughness variation of the borides formed on the HCCIs Boriding temperature (°C)
Boriding time (h)
Kc (MPa m1/2)
900
6 8 6 8
4.18 3.61 3.13 2.85
950 Fig. 6. Microhardness profile of the HCCI borided at 950 °C for 8 h.
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Table 2 Abrasive wear results of the borided HCCIs treated at different parameters Boriding temperature (°C)
Boriding time (h)
Wear rate (mg km− 1)
Relative wear resistance/ε
900
2 4 6 8 2 4 6 8
91.08 25.80 21.66 19.11 25.80 22.93 11.78 10.51
1.58 5.58 6.65 7.53 5.57 6.28 12.22 13.70
950
toughness of borided layers formed on the HCCIs ranges between 2.85 and 4.18 MPa m1/2 (Table 1). According to the previous studies [9,14], it was found that longer boronizing time and higher temperature are responsible for the lower fracture toughness of the borides. It was also reported the fracture toughness of the borides depends strongly on the presence of alloying elements in the substrates, such as Cr element which has a negative effect on the fracture toughness of borides [15]. Based on the XRD results mentioned above, there are borides mixture (FeB, Fe2B or CrB) formed on the surface. However, each one will have its own fracture toughness value, but it is not known what kind of interactions exist between different borides, and the influence of each on fracture toughness of composite boride layer is hard to discern [4,16,17]. 3.4. Abrasive wear behavior Table 2 presents the abrasive wear results of the borided HCCIs treated at different parameters. According to the table, the relative wear resistance of the borided HCCIs was approximately 1.58– 13.70, which indicates that the borided samples show much more excellent wear resistance compared with the traditional heat treated one. Particularly, the specimen borided at 950 °C/8 h, with the wear rate of 10.51 mg km− 1, shows the best wear resistance. In addition ,we can come to a conclusion from the data (Table 2) that the relative wear resistance depends closely on the boriding temperature and the processing duration. In other words, higher temperature and longer process duration result in more excellent wear resistance. It is well known that hardness of boride layer plays an important role in improvement of the abrasive wear resistance. As shown in Fig. 8, the
Fig. 9. Worn surface morphology of the HCCI borided at 950 °C for 8 h.
relationship between the surface microhardness and the relative wear resistance of the borided samples also confirms that the relative wear resistance was improved with the hardness increasing. Evans et al. [18] also found that fracture toughness is another significant factor that effects wear resistance. However, the fracture toughness of the borides varies in a narrow range (2.85–4.18 MPa m1/2) in our study, which is always one order of magnitude less than that of steels(50–150 MPa m1/2). Therefore, it can be considered that the fracture toughness of the borides have little influence on the wear resistance. So the higher hardness of the boride layer is thought to be the dominant factor for its higher relative wear resistance. This is in agreement with the previous studies reported in Refs. [19–22]. Worn surface morphology could help to understand the wear performance and the wear mechanism of the sample. Fig. 9 shows wear tracks of the HCCI sample borided at 950 °C for 8 h. It can be seen that mild abrasive wear with only a few shallower and narrower grooves was exhibited, and no delaminated debris can be observed. Thus, the micro-grooving is considered to be the dominant abrasive mechanism. The hardness value of the counterface material Al2O3 which is about 1800–2200 HV is close to the hardness of the boride layer. So the sample surface is difficult to be penetrated into. However, it can be seen that several paralleled grooves go through the structural defects (e.g. porosity or cracks), marked with the rectangulars in Fig. 9. Generally, structural defects which are tended to form the ploughing source always results in lower hardness and stress concentration and can't be avoided in the boride layer [4,7]. Consequently, decreasing the structural defects is also help for improving the resistance. 4. Conclusions The results of the current study can be summarized as follows:
Fig. 8. Relative wear resistance versus surface microhardness (HV0.5).
1. It is possible to develop a useful non-oxide boride type layer on the surfaces of HCCIs by conventional boriding treatment, and a compact and relatively smooth morphology was observed. 2. The boride layer formed at 950 °C/8 h is constituted of the phases : FeB, Fe2B and CrB while the layer formed at 900 °C/8 h, the phase present is mainly Fe2B. 3. Surface microhardness increased with increasing boriding time and higher temperature, and the boride layers formed at 950 °C showed a higher microhardness value compared with that at 900 °C. 4. The fracture toughness of borided layers formed on the HCCIs ranges between 2.85 and 4.18 MPa m1/2. 5. Abrasive wear test results show that the relative wear resistance of the borided HCCIs was about 1.58–13.70 and micro-grooving was considered to be the dominant abrasive mechanism.
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References [1] A. Wiengmoon, T. Chairuangsri, A. Brown, R. Brydson, D.V. Edmonds, J.T.H. Pearce, Acta Mater. 53 (2005) 4143. [2] H.H. Liu, J. Wang, B.L. Shen, H.S. Yang, S.J. Gao, S.J. Huang, Mater. Des. 28 (2007) 1059. [3] L.R. Shen, K. Wang, J. Tie, H.H. Tong, Q.C. Chen, D.L. Tang, R.K.Y. Fu, P.K. Chu, Surf. Coat. Technol. 196 (2005) 349. [4] I. Ozbek, C. Bindal, Surf. Coat. Technol. 154 (2002) 14. [5] B. Selcuk, R. Ipek, M.B. Karamis, V. Kuzucu, J. Mater. Process. Technol. 103 (2000) 310. [6] I. Campos, G. Ramírez, U. Figueroa, J. Martínez, O. Morales, Appl. Surf. Sci. 253 (2007) 3469. [7] Sukru Taktak, Mater. Des. 28 (2007) 1836. [8] C. Martini, G. Palombarini, G. Poli, D. Prandstraller, Wear 256 (2004) 608. [9] U. Sen, S. Sen, Mater. Charact. 50 (2003) 261. [10] C. Meric, S. Sahin, B. Backir, N.S. Koksal, Mater. Des. 27 (2006) 751.
[11] V.I. Dybkov, W. Lengauer, K. Barmak, J. Alloys, Compd. 398 (2005) 113. [12] Ugur Sen, Saduman Sen, Sakip Koksal, Fevzi Yilmaz, Mater. Des. 26 (2005) 175. [13] Wenran Feng, Dianran Yan, Jining He, Guling Zhang, Guangliang Chen, Weichao Gu, Size Yan, Appl. Surf. Sci. 243 (2005) 204. [14] I. Uslu, H. Comert, M. Ipek, F.G. Celebi, O. Ozdemir, C. Bindal, Mater. Des. 28 (2007) 1819. [15] Cuma Bindal, A. Hikmet Üçisik, Surf. Coat. Technol. 122 (1999) 208. [16] A.H. Ucisik, C. Bindal, Surf. Coat. Technol. 94/95 (1997) 561. [17] C. Bindal, A.H. Ucisik, Surf. Coat. Technol. 122 (1999) 208. [18] A.G. Evans, D.B. Marshall, D.A. Rigney (Ed.), ASM, Metals Park, OH, 1981: 439. [19] H.H. Liu, J. Wang, B.L. Shen, H.S. Yang, S.J. Gao, S.J. Huang, J. Iron Steel Res. Int. 13 (2006) 43. [20] H.S. Yang, J. Wang, B.L. Shen, H.H. Liu, S.J. Gao, S.J. Huang, Wear 261 (2006) 1150. [21] Z.P. Sun, R.L. Zuo, C. Li, B.L. Shen, S.J. Gao, S.J. Huang, J. Iron Steel Res. Int. 12 (2005) 34. [22] Z.P. Sun, R.L. Zuo, C. Li, B.L. Shen, S.J. Gao, S.J. Huang, J. Univ. Sci. Technol. Beijing 11 (2004) 539.
Contents lists available at ScienceDirect
Surface & Coatings Technology j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / s u r f c o a t
Effect of boronizing temperature and time on microstructure and abrasion wear resistance of Cr12Mn2V2 high chromium cast iron Chunmin Li, Baoluo Shen ⁎, Guijiang Li, Chao Yang College of Materials Science and Engineering, Sichuan University, Chengdu 610064, China
A R T I C L E
I N F O
Article history: Received 16 January 2008 Accepted in revised form 12 June 2008 Available online 27 June 2008 Keywords: High chromium cast iron Boriding Microstructure Abrasion wear resistance
A B S T R A C T In this study, Cr12Mn2V2 high chromium cast iron (HCCI) was boronized at 900 °C and 950 °C for 2, 4, 6, and 8 h, respectively. The borided samples were investigated by X-ray diffraction (XRD), scanning electron microscopy (SEM), microhardness tester and ring-on-block wear tester in terms of the phase composition, microstructure and worn surface morphology, microhardness, fracture toughness and abrasive wear resistance. The boride layer thickness ranges from 8 to 33 µm. XRD studies show the boride layer formed at 950 °C/8 h consists of the phases FeB, Fe2B and CrB, while for the layer formed at 900 °C/8 h, consists of mainly the Fe2B phase. Abrasive wear test results show that the relative wear resistance of the borided HCCIs increases with increasing surface microhardness. © 2008 Elsevier B.V. All rights reserved.
1. Introduction High chromium cast iron is widely used for wear components in the mining, mineral and oil industry due to its high hardness (because of the presence of hard (Fe,Cr)7C3 carbide) [1–2]. However, the cylinder liner that is made of HCCIs in mud pumps is very susceptible to damages in oil drilling machines. Improving the lifetime of such component is an important task. In recent decades, research activities have focused on surface treatments such as boriding, chromium coating, ion implantation and so on [3]. Boriding is thought to be an important thermochemical treatment to enhance the surface hardness and wear resistance of ferrous and non-ferrous alloy components [4]. Solid-state boriding has some important advantages in terms of easy handling, the flexibility with respect to the composition of the powder, minimal equipment and low cost [5–7]. The aim of this paper is to carry out solid-state boriding of Cr12Mn2V2 cast iron and investigate the resulting microstructure and properties of the borided Cr12Mn2V2 high chromium cast iron using X-ray diffraction (XRD), scanning electron microscopy (SEM), microhardness tester and ring-on-block wear tester. 2. Methods and procedures 2.1. Substrate materials Cr12Mn2V2 hypoeutectic HCCI samples were used for this study, with the following composition (wt.%):2.76%C, 12.62%Cr, 0.90%Si,
⁎ Corresponding author. Tel.: +86 13688456619; fax: +86 02885402231. E-mail address: [email protected] (B. Shen). 0257-8972/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2008.06.170
1.98%Mn, 1.6%V, 0.072%P, 0.036%S, and balance Fe. The HCCI was cast from 1500 °C as 80-mm-diameter balls into a chilled mould. Then the balls were cut into a size of 10 mm × 10 mm × 10 mm samples for the boriding treatments. Before the boriding treatment, all the samples were ground with 600 grid emery papers to get the final surface finish and then ultrasonically cleaned in acetone. 2.2. Treatments Boriding treatment was carried out in a sealed stainless steel container located in an electrical resistance furnace. The container was filled with the granular boriding medium that consisted of the boriding source B4C, the activator KBF4 and the diluent SiC. The test samples to be borided were placed in contact with the boriding medium. Then, boriding was performed at the constant atmosphere at 900 °C and 950 °C for 2, 4, 6, and 8 h, respectively, followed by furnace cooling. 2.3. Characterization The microstructure of the polished and etched cross-section of the specimens was observed by scanning electron microscopy (SEM, JSM5910LV). The thickness of boride layer was also measured using SEM. The Type Rigaku D/Max-rA X-ray diffraction apparatus was employed to examine the phase composition of the boride layer, with the radiation Cu-Kα, the tube voltage 42 kV and the tube current 110 mA. The microhardness profile of the boride layer was assessed by Type SHIMADZU microhardness tester with a load of 25 g. The surface microhardness and the fracture toughness of the boride layer were measured by Type MVC-1000A microhardness tester with the applied load 500 g and 300 g, respectively. The fracture toughness calculation was obtained by the equation Kc = 0.028(E/Hv)1/2 · P/c3/2[7], where E is
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Fig. 3. Cross-section view of the HCCI borided at 950 °C for 8 h. Fig. 1. Schematic representation of a Vickers indention in a coated material.
the Young's modulus of the boride layer (approximately 290 GPa), Hv is the layer hardness, P is the applied load, and c is the indentation half crack length as shown in Fig. 1. The abrasive wear test was carried out on a Type M200 ring-onblock wear tester. A schematic view of the tester is shown in Fig. 2. The test was performed at a constant speed of 200 rpm (equal to a liner velocity of 52.3 cm/s) under a load of 30 N. For the test duration, 10 min was selected .The counterface ring is a 50-mm-diameter 45#steel ring pasted with the 120#-alumina (Al2O3) cloth. Mass loss (M) of the sample was measured by TG328A photoelectric balance with accuracy of 0.1 mg. The wear rate was defined as mass loss per unit sliding distance (mg/km) and the relative wear resistance was defined as the ratio of wear rate of the uncoated cast iron (but held at 1000 °C for 0.5 h followed by air cooling) to that of borided cast iron. The worn surface morphology of the borided sample was also observed by SEM. 3. Results and discussion 3.1. Microstructure Fig. 3 shows the SEM image of the etched cross-section of the HCCI borided at 950 °C for 8 h. As can be seen, a compact and relatively smooth layer is formed on the surface and no obvious transition zone
Fig. 2. Schematic diagram of the abrasion tester.
is found. The morphology of the boride layers depends fundamentally on the chemical composition of the substrate. It is reported that sawtoothed layers are formed on the low-alloy steels while the layers are flat in the case of high-alloy steels [7,8]. As substrates, the Cr12Mn2V2 HCCI has a higher content of alloying elements, such as Cr, Mn and V, among which, Cr and V help with the formation of the FeB phase but inhibit the growth of saw-toothed layer. In the treating process, these atoms tend to concentrate at the tips of the boride columns, reducing considerably the active boron flux in this zone. Consequently, the columnarity was decreased at the interfaces [6]. In addition, the interface between the boride layer and the substrate is not obviously distinguished. This is probably due to the combination of borides and carbides. Fig. 4 shows the boride layer thickness of the HCCIs as a function of time. It can be observed that the thickness is parabolically increasing with time and its thickness value is in the range of 8–33 µm. As it is well known, thickness of the boride layer is closely related with the process temperature, treatment period and chemical composition of substrate [9]. Due to the fact that the process temperature plays a much more important role in the layer thickness increase than the treatment duration [7], the boride layers formed at 950 °C show larger thickness than that at 900 °C as shown in Fig. 4. On the other hand, it was found that the boride layer thickness of the Cr12Mn2V2 sample is much
Fig. 4. Boride layer thickness of the borided HCCIs.
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Fig. 7. Surface microhardness (HV0.5) of the HCCIs treated at different borided conditions. Fig. 5. XRD patterns of the borided HCCIs.
smaller than that of the other low and medium carbon plain carbon steels including low alloyed steels [6,9]. This phenomenon can be explained as follows: (1) Higher carbon content in the HCCI: since the carbon is not soluble in iron borides, it can be concentrated beneath the boride layer while diffusing into the substrate. So the boride diffusion was prohibited and the formation of the boride layer was hindered [10]. (2) The large amount of Cr and V in the HCCI substrates can reduce the active boron diffusivity by entering the iron boride lattice, as a results, the boride layer thickness tend to decrease [7]. 3.2. XRD analysis Fig. 5 reveals the XRD patterns of borided samples. As shown in Fig. 5, it can be observed that the diffraction peaks of the Fe2B produced on the surface of the 900 °C/8 h boronizing sample and 950 °C/ 8 h boronizing sample are high and sharp, however, the diffraction peaks of the FeB and CrB formed at 900 °C/8 h become weak or disappeared in comparison to that formed at 950 °C/8 h. This suggests that lower process temperature is favorable to form a single phase Fe2B for the HCCI. The reason for this is most probably attributed to the lower boron potential in the boriding process. Though there is 1.6% V
in the HCCI, no vanadium borides (e.g. VB) were observed from the results of all the XRD patterns. This is probably because vanadium borides are too little to be detected. 3.3. Microhardness and fracture toughness Fig. 6 shows the microhardness profile measured on the crosssection of the HCCI borided at 950 °C/8 h. It can be seen that the average value of microhardness of the borides reaches 1566 HV0.025, which is much greater than that of the substrate due to the present FeB, CrB and Fe2B phases in the boride layer. Moreover, it can also be observed that the microhardness fluctuated within a narrow range .This can be attributed to the presence of structural defects (e.g. porosity, cracks, etc.) and the different type of the borides [4]. Fig. 7 shows the surface microhardness(HV0.5) of the HCCIs treated at different borided conditions. As can be seen, the surface microhardness of the borided samples varies between 719 HV0.5 and 1650 HV0.5. It increases with boriding time extending and temperatures increasing. The FeB phase was found to be much harder than Fe2B phase [8,11], and Ugur Sen also showed [12] that the Vickers microhardness of the FeB phase varied between 1920 and 2140 HV0.1, the inner Fe2B phase ranged from 1160 to 1920 HV0.1 in the borides formed on the ductile iron substrate, so the boride layers formed at 950 °C show an obvious increase in hardness compared with those at 900 °C due to the FeB phase formation. Moreover, as a result of the large load (500 g) applied, borides formed on the HCCIs show an obvious phenomenon of indentation size effect (ISE, which usually involves an apparent decrease in the microhardness with the increasing applied test load) [13].Therefore, both the hardness of the single phase Fe2B and the indentation size effect are the dominant reason for the lower surface microhardness of the borides formed at 900 °C. Indentation fracture toughness tests are performed for the samples borided at 900 °C and 950 °C for 6 h and 8 h. The fracture
Table 1 Fracture toughness variation of the borides formed on the HCCIs Boriding temperature (°C)
Boriding time (h)
Kc (MPa m1/2)
900
6 8 6 8
4.18 3.61 3.13 2.85
950 Fig. 6. Microhardness profile of the HCCI borided at 950 °C for 8 h.
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Table 2 Abrasive wear results of the borided HCCIs treated at different parameters Boriding temperature (°C)
Boriding time (h)
Wear rate (mg km− 1)
Relative wear resistance/ε
900
2 4 6 8 2 4 6 8
91.08 25.80 21.66 19.11 25.80 22.93 11.78 10.51
1.58 5.58 6.65 7.53 5.57 6.28 12.22 13.70
950
toughness of borided layers formed on the HCCIs ranges between 2.85 and 4.18 MPa m1/2 (Table 1). According to the previous studies [9,14], it was found that longer boronizing time and higher temperature are responsible for the lower fracture toughness of the borides. It was also reported the fracture toughness of the borides depends strongly on the presence of alloying elements in the substrates, such as Cr element which has a negative effect on the fracture toughness of borides [15]. Based on the XRD results mentioned above, there are borides mixture (FeB, Fe2B or CrB) formed on the surface. However, each one will have its own fracture toughness value, but it is not known what kind of interactions exist between different borides, and the influence of each on fracture toughness of composite boride layer is hard to discern [4,16,17]. 3.4. Abrasive wear behavior Table 2 presents the abrasive wear results of the borided HCCIs treated at different parameters. According to the table, the relative wear resistance of the borided HCCIs was approximately 1.58– 13.70, which indicates that the borided samples show much more excellent wear resistance compared with the traditional heat treated one. Particularly, the specimen borided at 950 °C/8 h, with the wear rate of 10.51 mg km− 1, shows the best wear resistance. In addition ,we can come to a conclusion from the data (Table 2) that the relative wear resistance depends closely on the boriding temperature and the processing duration. In other words, higher temperature and longer process duration result in more excellent wear resistance. It is well known that hardness of boride layer plays an important role in improvement of the abrasive wear resistance. As shown in Fig. 8, the
Fig. 9. Worn surface morphology of the HCCI borided at 950 °C for 8 h.
relationship between the surface microhardness and the relative wear resistance of the borided samples also confirms that the relative wear resistance was improved with the hardness increasing. Evans et al. [18] also found that fracture toughness is another significant factor that effects wear resistance. However, the fracture toughness of the borides varies in a narrow range (2.85–4.18 MPa m1/2) in our study, which is always one order of magnitude less than that of steels(50–150 MPa m1/2). Therefore, it can be considered that the fracture toughness of the borides have little influence on the wear resistance. So the higher hardness of the boride layer is thought to be the dominant factor for its higher relative wear resistance. This is in agreement with the previous studies reported in Refs. [19–22]. Worn surface morphology could help to understand the wear performance and the wear mechanism of the sample. Fig. 9 shows wear tracks of the HCCI sample borided at 950 °C for 8 h. It can be seen that mild abrasive wear with only a few shallower and narrower grooves was exhibited, and no delaminated debris can be observed. Thus, the micro-grooving is considered to be the dominant abrasive mechanism. The hardness value of the counterface material Al2O3 which is about 1800–2200 HV is close to the hardness of the boride layer. So the sample surface is difficult to be penetrated into. However, it can be seen that several paralleled grooves go through the structural defects (e.g. porosity or cracks), marked with the rectangulars in Fig. 9. Generally, structural defects which are tended to form the ploughing source always results in lower hardness and stress concentration and can't be avoided in the boride layer [4,7]. Consequently, decreasing the structural defects is also help for improving the resistance. 4. Conclusions The results of the current study can be summarized as follows:
Fig. 8. Relative wear resistance versus surface microhardness (HV0.5).
1. It is possible to develop a useful non-oxide boride type layer on the surfaces of HCCIs by conventional boriding treatment, and a compact and relatively smooth morphology was observed. 2. The boride layer formed at 950 °C/8 h is constituted of the phases : FeB, Fe2B and CrB while the layer formed at 900 °C/8 h, the phase present is mainly Fe2B. 3. Surface microhardness increased with increasing boriding time and higher temperature, and the boride layers formed at 950 °C showed a higher microhardness value compared with that at 900 °C. 4. The fracture toughness of borided layers formed on the HCCIs ranges between 2.85 and 4.18 MPa m1/2. 5. Abrasive wear test results show that the relative wear resistance of the borided HCCIs was about 1.58–13.70 and micro-grooving was considered to be the dominant abrasive mechanism.
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