Effect of molybdenum on the microstructure and wear resistance of Fe-based hardfacing coatings

Materials Science and Engineering A 489 (2008) 193–200

Effect of molybdenum on the microstructure and wear resistance of Fe-based hardfacing coatings X.H. Wang a,∗ , F. Han b , X.M. Liu a , S.Y. Qu a , Z.D. Zou a a

b

School of Materials Science and Engineering, Shandong University, Jinan 250061, China Department of Mechanical and Electrical Engineering, College of Weifang, Weifang 261021, China Received 1 July 2007; received in revised form 7 December 2007; accepted 7 December 2007

Abstract Fe-based hardfacing alloys containing molybdenum compound have been deposited on AISI 1020 steel substrates by shield manual arc welding (SMAW) process. The effect of Mo on the microstructure and wear resistance of the Fe-based hardfacing alloys were investigated by means of X-ray diffraction, optical microscopy, scanning electron microscopy (SEM) and electron probe microanalysis, as well as wear test. The results indicated that cuboidal and rod-type complex carbides were synthesized in the lath martensite matrix. The fraction of carbides in hardfacing layer increased with an increasing of Mo content. The hardfacing layer with good cracking resistance and wear resistance could be obtained when the amounts of Fe–Mo was controlled within a range of 3–4 wt.%. The improvement of hardness and wear resistance of the hardfacing layers attributed to the formation of Mo2 C carbide and the solution strengthening of Mo. © 2007 Elsevier B.V. All rights reserved. Keywords: Microstructure; Hardfacing; Wear properties; Molybdenum

1. Introduction Wear of machine parts is one of the most common problems in engineering applications. To increase the wear resistance of those parts, Fe-based or nonferrous alloys with excellent wear properties have been researched and developed [1–5]. Nonferrous alloys, such as Co-based or Ni-based hardfacing alloys, are usually used for high-temperature application, they offer the overall combination of high corrosion or heat resistance coupled with wear resistance [2]. However, nonferrous alloys are more expensive than Fe-based hardfacing alloys. Fe-based alloys with relatively high concentrations of Cr and C have been widely used as hardfacing materials in applications. It has revealed that the formation of microstructures composed of ␣-ferrite and complex carbides, such as M7 C3 and M23 C6 [6]. But the carbides of Cr tend to precipitate along grain boundaries; it resulted in deterioration in toughness of the deposit coatings [7]. To avoid form the crack of the deposit coatings, a complicated welding process, such as pre- or post-welding heat treatment, should be employed during welding [8].

∗

Corresponding author. Tel.: +86 531 88392208; fax: +86 531 82616431. E-mail address: [email protected] (X.H. Wang).

0921-5093/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2007.12.020

In general, the excellent wear resistance results from high volume fraction of carbides and the toughness of the matrix [6]. To improve the toughness of the matrix, some literatures [7,9–11] reveals that strong carbide forming elements combined with carbon to form granular carbides and decrease the carbon content of the matrix, which is favorable for improving the toughness of the deposit coatings. They tried to fabricate carbides via reaction of Ti, V, Nb metallic elements and carbon through welding process. In our previous study, TiC–VC carbides reinforced Febased hardfacing layers have been fabricated by shield manual arc welding (SMAW) process [12]. It showed that TiC and VC carbides can be formed via metallurgical reaction of Fe–Ti, Fe–V and graphite. The macro-hardness of the deposited metals was over 60 HRC. However, due to the lower hardness of the matrix, it was found that there occurred selective wear of the matrix and formed voids at carbide/matrix interface, which induced fall-offs of carbide. Recently, some researches have proved that refractory metal of molybdenum is a solid-solution hardening element. The Mo can also contribute to the strength via precipitation hardening by forming MC or M2 C carbides [13–15]. The first objective of the present investigation is tried to increase the hardness of the matrix through molybdenum element which was added in the coating of the electrode. Second, the objective is to study the effect of molybdenum element on

194

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

Table 1 Composition (wt.%) of hardfacing electrode Samples

Fe–Ti

Fe–V

Fe–Mo

Graphite

CaCO3

TiO2

CaF2

S1 S2 S3 S4 S5

15 15 15 15 15

12 12 12 12 12

0 2 3 4 5

10 10 10 10 10

10 10 10 10 10

30 30 30 30 30

12 12 12 12 12

the microstructure and wear properties of Fe-based hardfacing alloys. 2. Experimental details Five types of electrodes, which coated different measures of powders of ferrotitanium (Fe–Ti), ferrovanadium (Fe–V), ferromolybdenum (Fe–Mo), and graphite (purity 99.5%), etc., were used as hardfacing alloys in the coating in the SMAW process. Composition of each hardfacing electrode is given in Table 1. The core wire of the electrode is H08A with diameter of 4.0 mm. The outside diameter of the coated electrode is about 6.82 mm. The main compositions of H08A, Fe–Ti, Fe–V and Fe–Mo alloys are listed in Table 2. The hardfacing was done on AISI 1020 low carbon steel substrate by means of SMAW process under direct current with a reverse polarity. The substrate registered a hardness value of about 200 HV0.2 . Before welding, the substrates were air blasted to remove stains and other contaminants from their surface, and then they were cleaned with acetone. The deposition was carried out in flat position; the current and travel speed were fixed in all the tests and no buffer layers were used. In order to reduce the effect of dilution, hardfacing bead was deposited four layers; the total thickness of the hardfacing layer was about 10 mm. The welding parameters were as follows: welding current 170–180 A, arc voltage 20–25 V and travel speed 25–28 cm/min. To assure the dryness of the electrode, it was baked in the furnace at 250 ◦ C for 2 h before welding. Samples were cut from the deposited layers using an abrasive cutting machine under water-cooling condition. All samples were etched with a solution of 3% nital. Microstructure was observed with a scanning electron microscopy (SEM) (Molel JSM-6380LA, Japan) and transmission electron microscope (TEM) (Model H-800, Japan) with energy-dispersive spectrometer (EDS). The overall composition was determined by X-ray fluorescence analysis; micro-zone composition and elementary line distribution were investigated using electron probe

microanalyzer (EPMA) (Model JXA-8800R, Japan). A type of D/Max-Rc X-ray diffractometer with Cu K␣ radiation operated at 60 kV and 40 mA was used to analyze phase structure. The macro-hardness of the hardfacing layer was measured by means of the average measurements taken from the top surface of the deposited layer using a Rockwell hardness tester (C-scale). The volume fraction of carbide particles in top surface of the deposited layer was measured by a computerized image analyzer. A block-on-ring wear testing was carried out without lubrication at room temperature using a friction and wear tester (Model MM200, China). The ring material of the wear couple was W18Cr4V with macro-hardness 62 HRC. The outer radius of the circular test ring is 40 mm, and the width is 10 mm. The tested specimens were machined to block with size of 35 mm × 10 mm × 10 mm. The wear conditions were a normal load of 98 N, a sliding speed of 0.84 m s−1 and a sliding distance of 1260 m. The wear volume was calculated using the following formula [16]: ⎧ ⎫    ⎨ π 2 b b ⎬ b V =w r2 − r 2 sin−1 − (1) ⎩ 180 2r 2 4⎭ where w is width of the specimens (mm), b is width of the wear track (mm) and r is out radius of wear ring (mm). The wear volume loss was calculated after a certain time interval. The coefficient of friction was calculated according to the following formula [17]: μ=

T rP

(2)

where μ represents the friction coefficient, T stands for the average frictional torque (N mm), P is the normal load (N) applied to wear specimens, and r is the outer radius of the ring (mm). The friction coefficient of the hardfacing layer was recorded in real time by a computerized data acquisition system equipped with an analog/digital converter. 3. Results and discussion 3.1. Microstructure characteristics The SEM morphology of S4-specimen is shown in Fig. 1. It can be seen that there are a number of small white particles with sizes about 1–3 ␮m embedded in the matrix. These white particles are irregularly cuboidal form with a dispersed distribution in the matrix. Although some areas show more particles than

Table 2 Composition (wt.%) of bare electrode and alloy Materials

C

Mn

Ti

V

Mo

Al

Si

S

P

Fe

H08A Fe–Ti Fe–V Fe–Mo

≤0.10 <0.10 <0.75 <0.15

0.30–0.55 – – –

– 35.0–45.0 – –

– – 45.0–55.0 –

– – – 65.0–75.0

– <8.0 – –

≤0.3 <3.0 <2.5 <4.0

≤0.03 <0.03 <0.10 <0.10

≤0.03 <0.03 <0.05 <0.05

Balance Balance Balance Balance

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

195

Fig. 1. SEM microstructural feature of the hardfacing layer for S4-specimen.

others, the distribution of particles is uniform in the matrix in general. Fig. 2(a) and (b) shows the X-ray diffraction spectrum of S4 and S1 specimens of the hardfacing layer, respectively. It indicated that the hardfacing layer is mainly composed of TiC, VC, Mo2 C, Fe3 C and ␣-Fe, which mean that the TiC, VC and Mo2 C carbides were formed during welding procedure. In addition, no remaining graphite, Fe–Ti, Fe–V or Fe–Mo was identified in the X-ray diffraction pattern, which implies that the synthesizing reactions of TiC, VC and Mo2 C are almost complete. According to X-ray analysis results, those white particles shown in Fig. 1 should be carbide. Fig. 3 shows the elements distribution based on the results of EPMA analysis for elements[C], [Ti], [V] and [Mo] along line A–B in the top surface of the hardfacing layer. Scanning line for composition analysis passes through three carbide particles. It can be seen from Fig. 3(b) that the intensities of Ti, V, Mo and C become higher when scanning line meets carbides. It indicates that the carbides are the complex carbides particles. At the same time, the further analyzing results for different regions of the matrix showed there molybdenum elements distributed

Fig. 2. X-ray diffraction spectrum of the hardfacing layer: (a) S4-specimen and (b) S1-specimen.

Fig. 3. EPMA analysis of TiC–VC–Mo2 C carbide particles for S4-specimen: (a) morphology of carbides; (b) line scan of Ti, V, Mo and C along line A–B.

into matrix, which indicates that part of Mo was dissolved in the matrix. A further investigation by TEM revealed that bulk carbides are precipitated in the matrix (as shown in Fig. 4). The corresponding electron diffraction pattern of mark (A) is shown in Fig. 4(b), which reveals the presence of TiC–Mo2 C carbide in the hardfacing layer. Fig. 4(c) shows the energy dispersive spectrograph (EDS) result of the carbide. It can be seen that the carbide contains mainly C, Ti, V and Mo elements, which also indicates that carbide is complex carbide containing at least two kinds of metallic elements. Fig. 4(d) and (e) shows the morphology of the matrix and corresponding electron diffraction pattern, respectively. Those indicate that the microstructure of the matrix is lath martensite. In general, lath martensite has higher hardness and better plasticity. Carbides embedded in the lath martensite matrix can help to prevent debonding between carbide and matrix. Therefore, hardfacing layer con-

196

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

Fig. 4. TEM image of the hardfacing layer for (a) image of carbides, (b) electron diffraction pattern of TiC–Mo2 C for marked A carbide, (c) EDS for marked A carbide, (d) image of lath martensite matrix, (e) electron pattern of matrix, and (f) EDS for matrix.

taining hard particles (carbide) and relative soft matrix (lath martensite) assists in reducing the crack. The EDS result of the matrix was shown in Fig. 4(f). It can be seen that some of Mo elements dissolved in the matrix, which indicates that Mo element not only forms Mo2 C carbides but also contributes to the strengthen the matrix via effect of solid-solution strengthening. The wear volume loss of S4-specimen as a function of sliding distance at a normal load of 98 N was shown in Fig. 5(a). Compared with the hardfacing layer produced by Mo-free hardfacing alloy, it can be seen that the hardfacing layer produced by hardfacing alloy with Mo possesses high wear resistance. For hardfacing alloy with Mo, besides formation of TiC or VC carbides, Mo element formed Mo2 C carbide, which contributes to the strength via precipitation hardening. Meanwhile, the matrix was strengthened by Mo element via solid solution. According to the test result of the hardness, the average microhardness of the matrix of the S4-specimen is about 789.4 HV0.2 , but the microhardness of the matrix of S1-specimen is about 609.6 HV0.2 . In general, hardness of the matrix is higher, the matrix can hold carbides effectively and inhibit the fall-off

carbides effectively, thereby leading to homogeneous wear of the matrix and carbides. The friction coefficient of hardfacing alloy with Mo is less than that of the Mo-free hardfacing alloy, as shown in Fig. 5(b). It can be seen that with the increasing of the distance, the friction coefficient of Mo-free hardfacing alloy undergoes a slight increase, but the hardfacing alloy with Mo varies slightly. In addition, for hardfacing alloy with Mo, it is found that there exist few of fluctuation in the values of the friction coefficient. A possible explanation for such fluctuations in the friction coefficient values has been suggested on the basis of a ’stick and slip’ mechanism [18]. As the asperities adhere during the wear, the moving parts stick, leading to a high friction values. The junction ruptures under the applied load the friction tends to zero. 3.2. Effect of Fe–Mo alloy on the microstructure and properties of hardfacing layer The overall chemical composition for each of the specimens is listed in Table 3.These data were measured at the top sur-

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

197

Table 3 Chemical composition (wt.%) of hardfacing layers Samples

C

Mn

Ti

V

Mo

Al

Si

S

P

Fe

S1 S2 S3 S4 S5

1.21 1.36 1.38 1.35 1.37

0.55 0.62 0.56 0.59 0.58

1.38 1.41 1.44 1.42 1.41

1.90 1.85 1.84 1.87 1.89

0.004 0.62 0.78 0.88 0.94

0.042 0.046 0.042 0.037 0.040

0.49 0.59 0.46 0.58 0.31

0.005 0.0035 0.004 0.003 0.005

0.023 0.021 0.018 0.026 0.022

Balance Balance Balance Balance Balance

face of the hardfacing layers. It can be seen that Mo content in the hardfacing layer increased with an increasing of Fe–Mo in the coating of the electrode. This resulted in a lower affinity of Mo for oxygen compared to the Ti and V elements. Thus, most of Mo could be transferred to the hardfacing layers. Fig. 6(a–e) reveals the effect of Fe–Mo on the microstructure of the hardfacing layers. From Fig. 6(a), it can be seen that there is a few carbides in the Mo-free hardfacing layer. It can also be noticed that the number of carbide and size of carbides are less than that of the hardfacing layer with Mo, as shown in Fig. 6(b–e). For hardfacing layer with Mo, cubic carbides and fine rod-shaped carbides can be observed in the composite coatings while the content of Fe–Mo is less than 5%. But for the S5-specimen that the content of Fe–Mo is over 6 wt.%, some network carbides are generally located along cell boundaries (as shown in Fig. 6(e)). This increased

Fig. 5. Wear properties of the hardfacing layer for S1 and S4 specimens: (a) wear volume loss and (b) friction coefficient.

the crack sensitivity of the hardfacing layer, as shown in Fig. 7. The effect of Fe–Mo on the hardness and volume fraction of carbide of the hardfacing layers was shown in Fig. 8. It can be seen that volume fraction of carbide and the hardness of the hardfacing layer gradually increased with an increasing of Fe–Mo content. The average hardness value was increased from 61.3 to 64.2 HRC when the content of Fe–Mo becomes higher. One reason of increasing hardness is attributed to the amount of carbides increased. On the other hand, more Mo elements dissolved in the matrix and enhanced the solution strengthening effect. Therefore, the hardness of hardfacing layers could be improved with the increasing of Mo content. Fig. 9 shows the effect of Fe–Mo on the wear volume loss of the hardfacing layers. It shows that the wear volume loss of the Fe-based alloy decreases with an increase in content of Mo. The increasing in the volume fraction of carbides leads to a corresponding increase in the hardness and wear resistance of the hardfacing layer. In addition, Mo dissolved in the Febased matrix could improve the bonding of the carbides and the matrix. It is reported for interred cermets that Mo reacts with TiC particles to form a case of (Ti, Mo)C1−x that surrounds the TiC core of each particle [19]. This mechanism tends to enhance the wettability of the carbide phase by the matrix and the result is strength enhancement. Therefore, wear occurs homogeneously and slowly due to very hard, fine complex carbides evenly distributed in the lath-type martensitic matrix. Fig. 10(a–e) shows the SEM micrographs of the worn surface of the S1–S5 specimens after wear test in air at normal load of 98 N and sliding distance 1260 m. Compared with the S2–S5 specimens, the worn surface of the Mo-free hardfacing layer (S1specimen) shows serious plastic deformation and deep groove, which results in the lower hardness and less volume fraction of carbide in the S1-specimen. However, for S2–S5 specimens, the hardfacing layers exhibit relatively smooth worn surface, and there is no indication of brittle failure or loose debris formation of carbides. It indicates that hardfacing layers with Mo alloy possess a much higher resistance to plastic erasing and removal of the edges of grooves. In addition, there are easily found some void formation at carbide/matrix interfaces and the fall-off of carbides in the S5-specimen. Compared to wear scar of S5-specimen, fall-offs of carbide of S2–S4 specimens are lower. Particularly in the S4-specimen, a smoothly worn surface is maintained with few-offs of carbides. It reveals that S3 and S4 specimens offer the best overall properties of wear and crack resistance.

198

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

Fig. 6. Effect of the amount of Fe–Mo on the microstructure of the hardfacing layer: (a) S1-specimen, (b) S2-specimen, (c) S3-specimen, (d) S4-specimen, and (e) S5-specimen.

Fig. 7. Crack of the carbide for S5-specimen.

Fig. 8. Effect of the amount of Fe–Mo on the volume fraction of carbides and the hardness of the hardfacing layer.

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

199

4. Conclusions

Fig. 9. Effect of the amount of Fe–Mo on the wear volume loss of the hardfacing layer.

(i) Complex carbides of TiC–VC–Mo2 C reinforced Fe-based hardfacing layers were prepared by SMAW process. Carbides were synthesized by means of metallurgical reaction. The forming carbides with cuboidal and rod-type were uniformly dispersed in the lath martensite matrix. (ii) The hardness and wear resistance of the hardfacing layer increases with an increasing of Fe–Mo content. When the amount of Fe–Mo is controlled within a range of 3–4%, hardfacing layer with better wear resistance could be obtained. However, when the amount of the Fe–Mo is over 5%, cracks were appeared in the hardfacing layer. (iii) Hardfacing alloy with Mo element alloy possess higher wear resistance than that of Mo-free hardfacing alloy. They exhibit relatively smooth worn surface, and possess high resistance to plastic deformation and scratch due to the effect of Mo solution strengthening.

Fig. 10. Wear scar of the hardfacing layer under a normal load of 98 N, a sliding speed of 0.84 m s−1 and sliding distance of 1260 m: (a) S1-specimen, (b) S2-specimen, (c) S3-specimen, (d) S4-specimen, and (e) S5-specimen.

200

X.H. Wang et al. / Materials Science and Engineering A 489 (2008) 193–200

Acknowledgements This research was supported by the Specialized Research Fund for the Doctoral Program of Shandong Province (2004BS04004) and Development Program of Shandong Province (06GG3203009). References [1] [2] [3] [4]

D.N. Noble, Met. Constr. (1985) 605–611. O.O. Zollinger, J.E. Beckham, C. Monroe, Weld. J. 77 (1998) 39–43. S. Chatterjee, T.K. Pal, J. Mater. Process. Technol. 173 (2006) 61–69. M.F. Buchely, J.C. Gutierrez, L.M. Le’on, A. Toro, Wear 259 (2005) 52– 61. [5] S.G. Sapate, A.V. Rama Rao, Wear 256 (2004) 774–786. [6] C. Fan, M.C. Chen, C.M. Chang, W. Wu, Surf. Coat. Technol. 201 (2006) 908–912. [7] Y.B. Zhang, D.Y. Ren, Mater. Sci. Technol. 19 (2003) 1029–1032.

[8] D.K. Dwivedi, Indian Foundry J. 47 (2001) 17–20. [9] B.D. Arnoldo, Int. J. Cast Met. Res. 13 (2001) 343–361. [10] R.K. Galgali, H.S. Ray, A.K. Chakrabarti, Mater. Sci. Technol. 15 (1999) 437–442. [11] S.L. Yang, X.Q. Lv, Z.D. Zou, S.N. Lou, Mater. Sci. Eng. A 438 (2006) 281–284. [12] X.H. Wang, M. Zhang, Z.D. Zou, S.Y. Qu, Mater. Sci. Technol. 22 (2006) 193–198. [13] M.M.A. Bepari, K.M. Shorowordi, J. Mater. Process. Technol. 155 (2004) 1972–1979. [14] M.X. Yao, J.B.C. Wu, Y. Xie, Mater. Sci. Eng. A 407 (2005) 234–244. [15] W. Wu, L.Y. Hwu, D.Y. Lin, J.L. Lee, Scripta Mater. 42 (2000) 1071– 1076. [16] H.L. Wang, H.H. Li, F.Y. Yan, Wear 258 (2004) 1562–1566. [17] H.B. Qiao, Q. Guo, A.G. Tian, G.L. Pan, L.B. Xu, Tribol. Int. 40 (2007) 105–110. [18] P. Lacey, A.A. Torrance, Wear 145 (1991) 367–383. [19] J.L. Elis, G.G. Goetzel, ASM Handbook, ASM International, Metal Park Ohio, OH, USA, 1993, pp. 978–1007.