ISIJ International, Vol. 61 (2021), No. 1

© 2021 ISIJ 434

* Corresponding author: E-mail: xuguang@wust.edu.cn

ISIJ International, Vol. 61 (2021), No. 1, pp. 434–441

https://doi.org/10.2355/isijinternational.ISIJINT-2020-327

1. Introduction

With the development and progress of the materials, the service scope of wear resistant steels (WRS) is moving towards large-scale, long-life and lightweight. Great atten-tions have been focused on improving the comprehensive properties of strength, hardness, toughness, plasticity and wear resistance.1–3)

Generally, martensite (M) is a dominant constituent in most kinds of WRS due to its good performance of wear resistance.4,5) Martensitic steels are usually treated by quenching and tempering to obtain the carbon-supersatu-rated martensitic phase with high hardness and strength. However, there are many problems such as cracking or warping which were caused by large internal stress after quenching in quenched steels.6,7) Furthermore, martensite is normally considered to be a brittle phase which is extremely detrimental to toughness. Therefore, many studies on mar-tensitic WRS focused on the improvement of toughness and plasticity without obviously compromising strength and hardness.8–10) In addition, bainite (B) has been proven to have good comprehensive performance of strength and toughness since it was discovered,11–15) and bainite WRS is developed by several researchers. For example, Kang et al. prepared a bainitic WRS with better comprehensive properties than the tempered martensitic WRS of the same grade.16) The wear resistance of a bainitic steel proposed by

Microstructure and Wear Properties of a Bainite/Martensite Multi-phase Wear Resistant Steel

Zishan YAO, Man LIU, Haijiang HU, Junyu TIAN and Guang XU*

The State Key Laboratory of Refractories and Metallurgy, The Collaborative Innovation Center for Advanced Steels of Ministry of Education, The Key Laboratory for Ferrous Metallurgy and Resources Utilization of Ministry of Education, Wuhan University of Science and Technology, Wuhan, 430081 China.

(Received on June 5, 2020; accepted on August 5, 2020)

The hardness, tensile properties, low temperature impact toughness and wear resistant performance in a low-carbon bainite/martensite multi-phase wear resistant steel were investigated. The results show that the multi-phase wear resistant steel with less amount of bainite exhibits better wear resistant performance at lower applied load, while the steel with more bainite displays better wear resistant performance at higher applied load. It is interesting to find that the wear rate varies with amount of baintie. Both the impact toughness and fracture elongation increase with the increase of the amount of bainite, while the hardness and strength decrease. The present work aims to provide guidance for the optimization of wear property, mechanical properties and microstructure of low carbon wear resistant steels.

KEY WORDS: bainite; martensite; multi-phase; wear property.

Li et al. was double compared to high manganese steel in the application of mining machinery.17) However, bainitic steels with better comprehensive properties normally require precise control of the heat treatment process and more addition of alloying elements, which makes it difficult to produce in industry.18,19)

In recent years, in order to simplify production pro-cess and enhance the comprehensive properties of steels, the multi-phase steels have been developed.20–22) Jha et al. reported that the best wear resistance was obtained in martensite/ferrite multi-phase steel compared to single mar-tensitic steel and ferrite/pearlite multi-phase steel with the same chemical composition.23) Tomita et al. reported that a small amount of bainite in martensitic steel could effectively improve the toughness with little decrease in strength.24) Normally dual-phase WRS normally has martensite as a main microstructure constituent. The research of the bainite/martensite dual-phase wear resistant steels (BMWRS) with bainite as a main constituent is rarely reported. At the same time, the wear behavior of BMWRS has not been clarified clearly.

In the present work, two different isothermal processes were utilized to produce different amounts of bainite in a low carbon alloy steel. And the comprehensive performance of the steel was analyzed, especially the wear resistance with different amounts of bainite. The results could provide theoretical guidance for the optimization of comprehensive properties and microstructure of low carbon BMWRS in different service conditions.

© 2021 The Iron and Steel Institute of Japan. This is an open access article under the terms of the Creative Commons Attribution license (https://creativecommons.org/licenses/by/4.0/).

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2. Experimental

The chemical composition of the investigated low car-bon BMWRS is shown in Table 1, in which alloying addi-tions were adopted to obtain desired mechanical and wear properties. Nickel (Ni) is normally considered to increase the low temperature impact toughness25,26) and molybde-num (Mo) increases the hardenability and the strength of steel, contributing to obtain bainitic microstructure.27,28) The experimental material was refined using a 50 Kg laboratory-scale induction furnace. Then, cast ingot was reheated and rolled to plates of 12 mm thick by a four-high mill. After the hot rolling process, the plate was fast-cooled to 500°C by laminar cooling and then air-cooled to room temperature.

Heat-treatment samples of 11 mm × 11 mm × 85 mm rectangle were taken from the hot-rolled plate along the rolling direction. Gleeble 3800 thermal-mechanical simu-lator was used to perform the heat-treatment experiments. Moreover, in order to determine the isothermal treatment temperature, the bainitic start temperature (Bs) and mar-tensitic start temperature (Ms) were calculated using the following empirical equations,29,30) and the results are given in Table 1.

Bs x x x x x

x x

Mn V Si Cr Ni

Mo W

� � � � � �� �630 45 40 35 30 20

25 15

Ms x x x

x x x

C Mn Cr

Ni Si Mo

� � � � � � �

� � � � � �

489 9 316 7 33 3 27 8

16 7 11 1

. . . .

. . xxW� �Figure 1 depicts the continuous cooling transformation

(CCT) curve of the tested steel plotted by JMat-pro 7.0 soft-ware, which provides guidance for the design of cooling rate before isothermal treatment. The experimental procedure is given in Fig. 2. Tested samples were austenitized at 900°C for 300 s followed by cooling to 340°C at a cooling rate of

20°C/s and isothermally kept for 200 s and 400 s, respec-tively. Afterwards, all the cases were quickly cooled at a rate of 50°C/s to room temperature. For simplicity, samples treated with isothermal times of 200 s and 400 s are termed as I200 and I400, respectively.

The metallographic samples were mounted and polished after heat-treatment experiments, and then etched by 4% nital. The microstructure observation of the as-perpared samples was carried out on a Zeiss metallographic opti-cal microscope (OM) and a Nova 400 scanning electron microscope (SEM). Moreover, the thin foil samples for transmission electron microscope (TEM) observation were prepared by twin-jet polishing technique using an electrolyte containing perchloric aicd (6%) and ethanol (94%), and the details of microstructure were observed by a JEM-2100F TEM. The amount of retained austenite (RA) was deter-mined using X-ray diffraction (XRD) using filtered Co Kα radiation with acceleration voltage of 40 kV and current of 40 mA.

Vickers hardness was measured by a MHVS-50Z hard-ness tester and the load of 500 g was chosen. The full-size Charpy U-notched impact specimens were machined based on ASTM standard, and the low temperature (−40°C) impact tests were performed using a JSCJ300-1 impact testing machine. After impact tests, the impact fracture surfaces were observed by SEM. Moreover, Instron 3380 electronic universal tensile machine was employed to test the tensile properties, and the stretching speed was set of 2 mm·min−1. Furthermore, the wear resistance experiments were conducted on a three-body impact abrasive wear resistance machine (MLD-10), which was used to simulate the working condition of mining machine components with certain impact energy. The schematic diagram of the wear resistance test device and the sample size as well as the counterface steel are given in Fig. 3. During the impact wear tests, the test samples were put in a clamping chuck which was connected with a hammer (10 kg) reciprocating motion, and the impact energy of 2 J and 4 J was chosen in this investigation, respectively. The frequency of ham-mer motion was set to be 180 beats per minute. And the different impact energy could be fulfilled by adjusting the height of the hammer. The abrasive of crushed silica sand particles (8–10 mesh, HV 950–1 250) were put into the

Table 1. Chemical composition, Bs and Ms of the tested steel (wt%).

C Si Mn P S Al Mo Ni Ms Bs

0.24 0.42 2.08 0.005 0.0034 0.023 0.20 1.50 313°C 488°C

Fig. 1. The CCT curve of the tested steel. Fig. 2. Schematic diagram of the experimental procedures.

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counter face steel and experimental steels at the rate of 1.4 kg/min, and the rotation speed of driving motor was 200 rolls per minute. GCr15 (58–61 HRC) was chosen for the martial of round counter face steel. It was austenitized at 840°C for 30 minutes and quenched to room tempera-ture, then heated to 180°C for 120 minutes for tempering treatment followed by furnace cooling. The samples were cleaned and polished by an ultrasonic cleaner before wear tests, and the loss of weight was measured after each test period of 600 s by an electronic balance. In addition, the SEM was used to observe the worn surface of each sample after wear test.

3. Results and Discussion

3.1. Thermal DilatationsThe thermal dilatometer was used to record the radial

expansions of the samples during thermal treatment. The temperature during the isothermal holding was well con-trolled, meaning that the dilatations were not affected by temperature fluctuation. Figures 4(a) and 4(b) illustrates the curves of dilatation versus temperature during the whole treatment cycle. As can be seen in Figs. 4(a) and 4(b), during the heating process, the transformation from α to γ started at 753°C and completed at about 800°C. Therefore, a heating temperature of 900°C (Fig. 2) could ensure suf-ficient austenitization of the tested steel. Moreover, it is a straight line during cooling process from 900°C to 340°C at a cooling rate of 20°C/s, indicating no high temperature transformation before isothermal holding treatment.

To clearly observe the variation trend of sample diam-eter during the isothermal holding and subsequent cooling process, the curves of dilatation after first fast cooling are enlarged in Fig. 4(c). The tangent method is used to confirm the starting and finishing temperatures of γ-α transforma-tion. From Fig. 4(c), the inflection points on the curves are marked as A, B, C, and D. Points A and B describe the starting and finishing temperatures of isothermal pro-cess, respectively. In addition, as shown in Table 1, the calculated Ms is 313°C. But the martensitic transformation occurred when the samples were cooled to 273°C (point C) and 251°C (point C1) for specimens I200 and I400, respec-tively. The difference between the calculated and measured Ms temperatures is attributed to the carbon enrichment in untransformed austenite due to the carbon diffusion dur-Fig. 3. Schematic of MLD-10 grinding abrasion test machine.

Fig. 4. Dilatation versus temperature curves of: (a) I200 and (b) I400; (c) dilatation-temperature curves and (d) dilata-tion-time curves during isothermal and subsequent cooling process. (Online version in color.)

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ing the bainitic transformation.31) The chemical stability of untransformed austenite increases with the increase of carbon content.32) Hence, the actual Ms of sample I200 decreased to 273°C. Furthermore, the carbon content in untransformed austenite increased with increasing amount of bainitic transformation at longer isothermal time of 400 s,33) leading to the further decrease of actual Ms in sample I400 (251°C).

Figure 4(d) reveals the curves of dilatation versus time during isothermal holding and subsequent cooling process. Obviously, as shown in Fig. 4(d), the dilatation-time curves of I200 and I400 maintain an increasing trend at the end of isothermal process, indicating the incomplete bainitic trans-formation even at the end of isothermal treatment of 400 s at 340°C. Thus, there was some untransformed austenite in both I200 and I400 specimens at the end of isothermal treat-ment. Part of residual austenite (untransformed austenite) transformed to martensite at a lower temperature and others were retained to room temperature as RA. Moreover, it can be observed that more bainite forms in specimen I400 due to the longer isothermal holding time.

3.2. MicrostructuresThe typical OM micrographs are shown in Fig. 5. It is

observed that the microstructure of two samples mainly consists of lath-like bainite and blocky martensite, and as expected, specimen I400 contains more bainite than speci-men I200. Compared with the large block characteristic in I200, the martensitic microstructure is refined in I400. The bainitic transformation is accompanied by carbon diffu-sion to adjacent austenite.34,35) When bainite transformation happens, carbon is rejected into adjacent untransformed austenite, resulting in chemical stabilization in the residual austenite. Carbon-rich residual austenite, therefore, would

transform to blocky martensite during the subsequent cool-ing process in I200. On the other hand, due to the more amount of bainite and less RA, the less amount of martensite is resultant in the sample of I400 (Fig. 5(b)). In addition, lath baintie could partition residual austenite into small blocks and more amount of baintie due to the longer transforma-tion time is likely to refine the multi-phase microstructure.

SEM microstructures are given in Fig. 6. As shown in Fig. 6(a) that the lath-like bainite in I200 parallelly distrib-utes in the grains and grows through the original austenite grain. Most of the large blocky martensite and/or martensite/austensite (M/A) islands are distributed between the bainite sheaves with different orientations in the prior austenite grain. This is because the original austenite grain was parti-tioned by the bainite sheaves during the isothermal holding and the untransformed blocky austenite either decomposed into martensite or partly retained to form M/A depending on its stability during the subsequent cooling process. In addi-tion, it is observed from Fig. 6(b) that the amount of lath bainite in I400 is obviously larger than that in I200, which is consistent with dilation results in Fig. 4. At the same time, the large blocky M/A disappears in I400.

Typical TEM micrographs of samples I200 and I400 are shown in Fig. 7, where the morphologies of B, M and RA are clearly presented. The martensite in TEM micrographs is darker than bainite due to the higher dislocation density. As shown in Fig. 7(a), film-like RA is observed between the bainitic laths. Moreover, compared to I200, the length of bainitic laths in I400 is obviously longer (Fig. 7(b)). Besides, intercept method was used to measure the width of bainitic laths based on the TEM micrographs.36) The width of bainitic laths of I200 and I400 is estimated to be 0.73 ± 0.15 μm and 0.66 ± 0.36 μm, respectively. The width of bainitic laths of specimens I200 and I400 are simi-lar because of the same isothermal temperature.

3.3. Quantitative AnalysisFigure 8 illustrates the different diffraction peaks of

specimens I200 and I400 by XRD tests. The angles and the integrated intensities of different diffraction peaks were measured by the software of High-Score Plus. According to the results in Fig. 8, the amount of RA were counted based on Eq. (3)

VG

I

I

i �� �

1

1 �

�

.............................. (3)

where the Vi in Eq. (3) represents the amount of RA for Fig. 5. The typical OM microstructures of: (a) I200 and (b) I400. (Online version in color.)

Fig. 6. The SEM microstructures of: (a) I200 and (b) I400. Fig. 7. The TEM microstructures: (a) I200 and (b) I400.

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each diffraction peak; the value of G is selected according to the study by Wang et al.,37) 2.5 for Iα (200)/Iγ (200), 1.2 for Iα (211)/Iγ (200); Iα and Iγ represent the corresponding integrated intensities of ferrite and austenite respectively. Furthermore, the carbon contents in RA were calculated by Eq. (4).38,39)

C�

��

�� �

��

���

�

���

h k l2 2 2

23 578 0 033

sin. / . ........... (4)

The values of angle of (200) austenite diffraction peaks are shown by θ, λ (1.78892 Å) is the wavelength of the Co Kα in XRD tests, and h, k and l are the Miller index of aus-tenite plane. The results are shown in Table 2. As shown in Table 2, the amount of RA increases from 3.12% to 6.64%, and the carbon content of RA increases from 0.279% to 0.664% with the increase of isothermal time from 200 s to 400 s, respectively. Because of the volume fraction of bainite increases with the increasing isothermal time, more carbon atoms are rejected into untransformed austenite, resulting in the larger chemical stability of untransformed austenite after the longer isothermal holding. Therefore, the volume fraction of RA and the content of carbon in RA in I400 are larger than that of I200.

The amount of bainite in I200 and I400 can be calculated based on the results of dilatation curves in Fig. 4. As shown in Fig. 4(c), for I200, the increment of dilatation from point A to B (labeled as X1) represents the amount of bainitic transformation. Secondly, the tangent line is plotted from the beginning point C of martensitic transformation, and then the vertical line is plotted from the end point D of mar-tensitic transformation. The difference between the ordinate of the intersection (Point E) of the two lines and the ordinate of point D (labeled as X2) reflects the amount of the mar-tensitic transformation. Thirdly, the amount of RA (Table 2) is labeled as X3. Finally, the phase percentage of bainite (labeled as X4) is calculated by X4 = X1/(X1 + X2)·(1 − X3). Based on this method, the bainite fractions in I200 and I400 are 68.72% and 82.06%, respectively, indicating that the amount of bainite increases by 19.41% with the increasing isothermal time from 200 s to 400 s. This is because the

bainitic transformation was not completely finished during isothermal holding of 200 s, and more bainite formed in specimen I400 during continuous isothermal holding.

3.4. Hardness and Low Temperature Impact Tough-ness

The hardness tests were performed on a HV-1000A micro-hardness tester after thermal simulation experiments. The Vickers hardness was measured for five times for each sample and the average value was obtained. The results are given in Table 3, in which the hardness of I200 is higher than I400. This is because the amount of martensite in I200 is larger than that in I400, resulting in a higher hardness in specimen I200. Therefore, longer isothermal time of 400 s has a negative effect on hardness due to more amount of bainite.

Additionally, the U-notch Charpy impact toughness experiments were performed at −40°C and the results are given in Table 3. Compared to I400, I200 shows worse low temperature impact toughness. This is attributed to the for-mation of more amount of large blocky martensite in I200 (Fig. 6), which is detrimental to impact toughness.40) The fracture surfaces were analyzed by SEM (Fig. 9), manifest-ing a typical ductile fracture with large number of ductile dimples in both I200 and I400. However, large dimples are surrounded by small dimples in the sample of I400 (Fig. 9(b)), demonstrating the better low temperature toughness of I400, while some river-like quasi-cleavages can be found in the fractograph of I200 (Fig. 9(a)). The micro-morphology of the fractures well matches with the results of the impact toughness tests in Table 3. Hence, longer isothermal time of 400 s has a positive effect on impact toughness.

Table 2. Results of RA and carbon content in RA.

Sample RA (%) Carbon content of RA (wt.%)

I200 3.12±1.35 0.279± 0.06

I400 6.64± 2.18 0.664± 0.09

Fig. 8. XRD patterns for specimens I200 and I400. (Online ver-sion in color.)

Fig. 9. Fracture surfaces of: (a) I200 and (b) I400.

Table 3. Hardness and low temperature (− 40°C) impact tough-ness of the tested steels.

Sample Hardness (HV) Low temperature impact toughness (J)

I200 464± 28 41.3± 3.5

I400 439±16 53.2± 2.7

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3.5. Tensile PropertiesFigure 10 depicts the curves of engineering stress-strain

of I200 and I400 and the tensile results are summarized in Table 4. It can be observed from Fig. 10 and Table 4 that the tensile strength (Rm), yield strength (Rp0.2) and product of tensile strength and elongation (PSE) decrease, while the fracture elongation (FE) increases with the increase of isothermal holding time. During the longer isothermal duration, the amount of RA is expected to be larger because of enhanced carbon diffusion, which results in the higher elongation due to more transformation-induced plasticity (TRIP) effect.41,42) At the same time, more amount of bainite also contributes to the increase of elongation. Compared with I400, the Rm and Rp0.2 of I200 increase by 11.65% and 4.86%, respectively. This is ascribed to more martens-ite transformation strengthening in microstructure of I200 (Fig. 6). Furthermore, the ratio of yield to tensile strength (Y/S) of I400 is larger than that of I200. Relevant studies have shown that the higher Y/S in steels represents better ability of deformation resistance.43) It is speculated that the deformation resistance of I400 is better than that of I200.

3.6. Wear Resistance and Worn SurfaceThe MLD-10 dynamic abrasive machine was used to

measure the wear resistance performance of specimens I200 and I400. During the wear process, the remaining weight of each sample was recorded every 600 s to count the loss of weight per unit time. The results of weight loss at dif-ferent wear times are presented in Fig. 11. It is shown that the weight loss of I400 is larger than that of I200 during whole wear resistance process at the applied load of 2 J. The impact wear morphologies observed by SEM are shown in Fig. 12. It reveals that the wear surface of I200 (Fig. 12(a)) with the load of 2 J is mainly dominated by plastic fatigue morphology, in which the area of ploughing/cutting

morphology and chiseling morphology are relatively small. Compared with I200 (Fig. 12(a)), the area of ploughing/cutting morphology obviously increases in I400 under the load of 2 J (Fig. 12(c)), and the plastic fatigue and abrasive embedded morphologies decrease. During the impact wear process under load of 2 J, the abrasive is easily embedded in the tested steel under the normal stress by the impact of the hammer in I400 specimen due to the lower hardness. After the abrasive is embedded into the tested steel, the steel surface is cut to form furrows under the shear stress of the rotation of the counterface steel. Therefore, the wear surface of I400 is mainly dominated by ploughing/cutting morphology. In addition, more amount of martensite leads to high hardness and strong resistance to abrasive cutting ability, and results in the less ploughing/cutting morpholo-gies in I200 (Fig. 12(a)). At the same time, the portion of plow-wrinkle of the extruded part in the experimental steel continuously underdoes deformation under repeated

Table 4. Results of tensile tests.

Sample Rp0.2 (MPa) Rm (MPa) FE (%) Y/S (%) PSE (GPa%)

I200 1 183± 28 1 390± 21 8.26± 0.4 0.85± 0.03 11.48± 0.81

I400 1 128± 33 1 245± 29 8.63± 0.2 0.90± 0.05 10.74± 0.62

Fig. 10. The engineering stress-strain curves.

Fig. 11. The weight loss of two steels at different wear times.

Fig. 12. The morphologies of wear surfaces: (a) I200 at 2 J; (b) I200 at 4 J; (c) I400 at 2 J; (d) I400 at 4 J.

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normal stress of impact. The hardness of this part gradually increases with the increase of plastic deformation, resulting in less weight loss. Hence, the morphology of wear surface of I200 is characterized by plastic fatigue morphology, and the weight loss of I200 is less than I400. In short, relatively smaller amount of bainite in BMWRS is conducive to the improvement of wear resistance under low impact load.

On the other hand, under impact load of 4 J, the weight loss of I400 is similar with I200 during the early 1 500 s. However, the weight loss of I200 significantly increases and exceeds that of I400 after 1 500 s. As can be seen in Figs. 12(b) and 12(d), the wear surface of both samples at load of 4 J is mainly dominated by ploughing/cutting morphol-ogy. It should be noted that except for ploughing/cutting morphologies, large fatigue spalling morphology can be observed in specimen I200 (Fig. 12(b)), which means larger weight loss. On the contrary, the appearance of fatigue spall-ing decreases in specimen I400 (Fig. 12(d)), meaning less weight loss. This is because the plastic deformation degree and the plastic fatigue area increase in I200 under repeated impact of abrasive at the impact load of 4 J. At the same time, with increasing wear time, the area of repeated plastic deformation is easy to fracture and spall due to the relatively lower low temperature impact toughness and worse fracture resistance, which results in the appearance of large spall-ing morphology (Fig. 12(b)) and leads to a large weight loss in I200. Furthermore, the deformation of wear sample increases with the increase of impact load during the wear process. At the same time, more TRIP effect at large defor-mation due to the higher amount of RA in I400 results in the increasing hardness and decreasing weight loss. There-fore, the weight loss of I200 exceeds I400 after 1 500 s. In a word, more amount of bainite in BMWRS is conducive to the improvement of wear resistance performance under large impact load.

Additionally, the wear rate is represented by the slope of weight loss curves every 600 s of wear time in Fig. 11. From Fig. 11, the wear rate of I400 is basically similar during whole wear resistance process at load of 2 J or 4 J, while the wear rate of I200 shows obviously increase after 2 400 s at load of 2 J and after 1 200 s at load of 4 J. It is observed in Figs. 12(a) and 12(b) that the area of ploughing/cutting morphology increases significantly, and the area of plastic fatigue morphology disappears basically, while the fatigue spalling and chiseling appear with the increasing load from 2 J to 4 J in I200. During the wear process, part of material in I200 extruded by the abrasive continuously underdoes plastic deformation, and gradually the extruded material reaches strain fatigue under repeated impact with the increasing wear time. And the toughness of this part of the material gradually decreases with the increase of plastic deformation. Eventually, this part of material is exfoliated due to the smaller impact toughness in I200, and results in the serious weight loss. Therefore, the wear rate of I200 increases at certain time during wear process.

In addition, the flection of wear rate in specimen I200 at 4 J is earlier than that at 2 J, which can be attributed to the following reasons. First of all, at impact load of 4 J, the abrasive is pressed deeper on the tested steel, resulting in the increase of the extruded part caused by the abrasive under shear stress. This leads to the increase of the plough-

ing/cutting morphology in I200 under impact load of 4 J. Secondly, the strain fatigue of part extruded by abrasive is more likely to occur due to the increase of the deformation degree under the large impact load. Thirdly, microcracks are relatively easy to form on the subsurface of the sample and more likely to extend to the surface under large impact load, which results in the appearance of fatigue spalling and chiseling in I200 under load of 4 J. Finally, the part of abrasive extruded in I200 is easier to be deformed and fractured under lager impact load due to the lower ability of deformation resistance and impact toughness. These factors contribute to the earlier increase in the wear rate of I200 (1 200 s) under large impact load.

On the other hand, the wear surfaces of I400 at load of 2 J and 4 J are mainly dominated by ploughing/cutting morphology, indicating that the weight loss of I400 mainly consists of abrasive cutting wear. In addition, the weight loss in I400 caused by abrasive cutting is larger due to the lower hardness and strength. Besides, the weight loss caused by strain fatigue and spalling is relatively less due to the better deformation resistance and fracture resistance of I400. Therefore, the wear rate of I400 is basically similar during whole wear resistance process at same loads. As mentioned above, the BMWRS with relative more amount of bainite displays a uniform wear rate under different impact loads.

4. Conclusions

Two isothermal procedures were designed to investigate the effect of bainite/martensite dual-phase microstructure on the hardness, tensile properties, low temperature impact toughness and wear property in a high-strength low carbon steel. The main conclusions are summarized as follows:

(1) The weight loss of steel with less bainite (I200) is smaller than that of steel with more bainite (I400) at small load of 2 J, indicating the better wear resistant performance of steel containing less bainite at low impact load for bainite/martensite dual-phase resistant steels, while steel with more bainite exhibits better wear resistant performance at large load of 4 J.

(2) The rate of wear loss of steel with less bainite increases after 2 400 s and 1 200 s during the wear process at load of 2 J and 4 J, respectively, while the rate of wear loss of steel with more bainite shows less fluctuation.

(3) The impact toughness, fracture elongation and ratio of yield to tensile strength increase, while the hardness, ten-sile strength, yield strength and product of tensile strength and elongation decrease with the increase of bainite amount in bainite/martensite dual-phase were resistant steels.

AcknowledgmentThe authors gratefully acknowledge the financial supports

from the National Natural Science Foundation of China (NSFC) (Nos. 51874216 and 51704217), the Key Project of Hebei Iron and Steel Group (HG2019313), the Major Projects of Technology Innovation of Hubei Province (No. 2017AAA116) and the Special Project of Technological Innovation of Wuhan (2018010402011187).

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