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Microstructures and properties of carbidic austempered ductile Iron containing Fe3C particles and superfine ausferrite
Materials and Design 186 (2020) 108363
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Microstructures and properties of carbidic austempered ductile Iron containing Fe3C particles and superfine ausferrite Yang Penghui a, Fu Hanguang a,⁎, Li Guolu b, Liu Jinhai b, Zhao Xuebo b a b
Key Laboratory of Advanced Functional Materials, Ministry of Education, School of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, PR China School of Materials Science and Engineering, Hebei University of Technology, Tianjin 300130, PR China
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• The impact toughness of CADI is significantly improved by super-high temperature pretreatment and austempering treatment. • The CADI obtained by S&A treatment has excellent wear resistance under high wear load. • With the rise of reheating temperature, the pearlite decreases gradually, and the precipitated particles are retained. • A superfine ausferrite matrix with a length of 1-3 μm and a thickness of 50 ± 10 nm is obtained. • The boundaries between bainitic ferrite and precipitated particles are coherent boundaries.
a r t i c l e
i n f o
Article history: Received 5 June 2019 Received in revised form 22 October 2019 Accepted 13 November 2019 Available online 14 November 2019 Keywords: CADI Super-high temperature pretreatment Superfine ausferrite Precipitates Impact toughness
a b s t r a c t A new type heat treatment process comprising super-high temperature pretreatment and austempering treatment (S&A treatment) was used to process carbidic austempered ductile iron (CADI). The results showed that the netlike eutectic carbides were significantly reduced after super-high temperature pretreatment. Due to the increase in supercooling during super-high temperature pretreatment cooling, the inter-laminar spacing of pearlite was reduced from 300 nm to 100 nm. When the ductile iron with superfine pearlite was reheated, a large amount of Fe3C particles were retained inside the prior austenite grains. After austempering, a superfine ausferrite matrix with a length of 1–3 μm and a thickness of 50 nm was obtained. The main reason for this refinement is the increased number of nucleation-sites for the austenite grains and the retardation-effect of Cr on their growth. TEM results indicated that a possible orientation relationship between bainitic ferrite with a cubic lattice and precipitated Fe3C particles with a orthorhombic lattice is (011)α-Fe//(210)Fe3C, and their boundaries are coherent boundaries. The CADI obtained by S&A treatment has an impact toughness 120% higher than the traditional CADI without sacrificing hardness, and has excellent wear resistance under high wear load. © 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
⁎ Corresponding author at: School of Materials Science and Engineering, Beijing University of Technology, Number 100, Pingle Garden, Chaoyang District, Beijing 100124, PR China. E-mail address: [email protected] (F. Hanguang).
1. Introduction Fracture, corrosion and wear are important failure modes of mechanical parts, which will cause huge economic losses. Therefore, the
https://doi.org/10.1016/j.matdes.2019.108363 0264-1275/© 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
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research on improving the wear resistance and corrosion resistance of materials has been highly valued by researchers. Austempered ductile iron (ADI) is widely used in motor-dom, mining, railway, agricultural machinery and other fields due to its high tensile strength (≤ 1600 MPa), good impact toughness (≥ 100 J/cm2), excellent fatigue resistance and corrosion resistance [1–4]. However, the wear resistance of ADI under high load is not ideal. To improve the wear resistance of ADI, researchers [5–8] invented a new wear-resistant and corrosion-resistant material called Carbidic Austempered Ductile Iron (CADI) on the basis of ADI. There were three main preparation processes for producing CADI: (1) Adding a small amount of alloying elements (such as Nb, V, Mo, Cr, etc.) into ADI to promote the precipitation of eutectic carbides [9–12]; (2) Setting chills in the mould to increase the cooling rate [13]; (3) Reducing the amount of graphitizing elements (in particular Si). Among them, the addition of alloying elements could shorten the crystallization temperature interval between the stable and the metastable eutectic reactions significantly, which promoted the non-equilibrium solidification and facilitated the formation of eutectic carbides. In addition, the addition of Cr could bring more carbides at lower production cost comparing with the same amount of other alloying elements [14]. However, the Cr contained eutectic carbides distributed at grain boundaries, which seriously destroys the continuity of the austenite matrix, resulting in the decrease of impact toughness [14]. Although the effects of alloying elements and cooling rate on hardness, impact toughness, tensile strength and microstructures of CADI have been reported in the past decades [9–14], the important mechanism for improving the toughness of CADI had not been provided. Until 2013, Sun et al. [15,16] proposed a method to improve the impact toughness of CADI using nano-oxides. As a result, when 0.06 wt% nanoceria was added into CADI containing 3.5C-2.85Si-0.75 wt% Cr, the impact toughness was increased by 33% due to the refinement of netlike eutectic carbides. Besides, the abrasive wear rate was reduced by 29%. In addition, the morphology and the size of eutectic carbides were changed by this method, which ignored the refinement of austenite matrix. It has been proved that fine bainitic plates have excellent comprehensive mechanical properties [17–21], such as nanostructured bainitic steel. These bainites are composed of very fine bainitic ferrite plates (about 50 nm) and retained austenite films. Due to the uniform microstructure and the low residual stresses, this bainitic steel has unique advantages in the application of large mechanical components. However, considering the phase transformation kinetics of bainite, the formation of these bainites usually requires long holding time (sometimes longer than several days) at the low temperature, or multi-pass rolling before bainitic transformation [22], which significantly reduces production efficiency. Liu et al. [23] approached the feasibility of adopting super-high austenitizing temperature (960–1000 °C) to improve the impact toughness of CADI. When adding 0.5 wt% Cr and adopting superhigh austenitizing temperature of 960–1000 °C, the impact toughness values (absorbed energy) of CADI specimens with the hardness of 39–50 HRC could reach 28–93 J/cm 2 . However, the grains became coarsened with the increase of austenitizing temperature, and the precipitates had not been found in CADI obtained by super-high austenitizing temperature. When adding 1.0 wt% Cr, the hardness of CADI had exceeded 50 HRC, and the toughness might be under 10 J/cm 2 . Therefore, a method that could greatly improve the impact toughness of CADI without sacrificing wear resistance had not been proposed to date. The present work aims to refine the eutectic carbides and the ausferrite matrix by an innovative heat treatment process including super-high temperature (closed to the melting point) pretreatment and austempering treatment (referred as S&A treatment). The proposed treatment process successfully improves the impact toughness and the wear resistance of CADI owing to the
replacement of the coarse eutectic carbides and coarse ausferrite matrix with lots of precipitated particles and superfine ausferrite, particularly, the wear resistance under medium and high wear load, and the research results are expected to accelerate the popularization and application of CADI, such as digger teeth, balls of ball mill, liners of ball mill, loose soil plow shovels, crusher and some automobile parts et al. 2. Experimental procedures 2.1. Material preparation and heat treatment The charge materials were melted using a 50 kg medium frequency induction furnace. 0.15 wt% pure aluminum was added to the melt to deoxidize the molten iron at 1480–1500 °C, and then transferred to a ladle containing 1.2 wt% inoculant and 1.5 wt% nodularizer. Spheroidization and inoculation were performed by pour-over treatment process, and inoculant and nodularizer were FeSi75 (75 wt% Si) and FeSiMg6Re2 (6Mg-2Re-40 wt% Si), respectively. After removing the slag, the melt was poured into sodium silicate-CO2 bonded sand molds, obtaining Y-block ingots according to ASTM A781/A 781-M95. The chemical composition of the ingots was measured by QualiSpark 750 Optical Emission Spectrometer, which was 3.72C-2.77Si-0.51Mn-0.99 wt% Cr. Samples with dimensions of 11 mm × 11 mm × 55 mm and 13 mm × 13 mm × 19 mm were cut from the bottom of Y- block ingots for impact test, microstructural analysis and wear test. Fig. 1a shows a traditional austempering treatment. The operation processes are as follows: The samples were heated to 880 °C for 120 min, and then instantly quenched into salt-bath of 50 wt% KNO350 wt% NaNO3 and isothermally held for 120 min at 300 °C, followed by air-cooling, called “Ordinary austempering treatment” below. Fig. 1b shows a super-high temperature treatment, Here, the sample was heated to 1100 °C for 10 min, cooled to 880 °C in the furnace, and held for 120 min, followed by isothermal quenching, called “Contrast test”. Fig. 1b also shows a water quenching treatment (dashed line), and the operation processes are as follows: The sample was heated to 1100 °C for 10 min, cooled to 880 °C in the furnace, and held for 120 min before water-quenched, called “Contrast WQ test”. To improve the mechanical properties of CADI and refine the ausferrite, a heat treatment process including super-high temperature pretreatment and austempering treatment used in the present work is shown in Fig. 1c (black lines), called “S&A treatment” below. The specific operation processes are as follows: Firstly, the sample was heated to 1100 °C for 10 min, then air cooled to room temperature. Secondly, the samples were reheated to 880 °C for 120 min, and then quenched into saltbath. Finally, isothermally held for 120 min at 300 °C, followed by aircooling. To more conveniently observe the precipitated particles, these samples were reheated to 720 °C, 760 °C, 800 °C, 840 °C, and 880 °C respectively after super-high temperature pretreatment with a duration of 120 min, followed by water-quenched. 2.2. Microstructural analysis The microstructures were investigated by an Olympus BX51 optical microscope (OM), a JSM-6510 scanning electron microscope (SEM), a Rigaku-Utima-IV X-ray diffraction (XRD), a JEM-2100F transmission electron microscopy (TEM) with an energy dispersive X-ray (EDX), an electron probe micro-analyzer (EPMA) with a wavelength dispersive X-ray (WDS). Cu-Kα was used to perform the radiation at 40 kV and 30 mA. The specimens were scanned in the angular 2θ range from 20° to 90° with a step size of 0.5°/min and a collection time of 10 s. TEM samples preparation processes were as follows: firstly, the samples with an area of 100 mm 2 and a thickness of 0.5 mm were cut by wire cutting; then mechanically ground by hand to a thickness of 30–50 μm without cracks and
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Fig. 1. Schematic diagram of heat treatments; (a) ordinary austempering treatment, (b) contrast test, (c) S&A treatment. SQ: salt quenching, AQ: air quenching, FC: furnace cooling, WQ: water quenching.
breakage; finally, thinning by plasma at 4.5 kV, 1.9 mA and rotating angle of 6°. The area of precipitated particles and eutectic carbides was calculated by Image-pro software. 20 fields of view were
randomly chosen, and the average value was calculated as the final result. The volume fraction and carbon content of retained austenite were estimated by XRD. The calculation formula of
Fig. 2. Microstructure of as-cast carbidic ductile iron: (a) optical microstructure (etched by 4% Nital), (b) SEM image (etched by 4% Nital), (c) Mapping distribution of Cr elements, (d) Bright-field TEM micrographs and corresponding selected area diffraction patterns (SADPs) of eutectic carbides.
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where ω γ is the volume fractions of austenite, ω α is the volume fractions of ferrite, I γ is the integral intensities of austenite, I α is the integral intensities of ferrite. K = 1.39. The carbon content of retained austenite is calculated by the following Eq. (2) [2]: aγ ¼ 0:3548 þ 0:0044 C γ
ð2Þ
where Cγ is the carbon content of retained austenite (wt%), aγ is the lattice constants of austenite (Å). 2.3. Mechanical performance test Fig. 3. XRD patterns of as-cast carbidic ductile iron, ordinary CADI and CADI of after S&A treatment: (a) as-cast, (b) ordinary CADI, (c) S&A treatment.
volume fraction is as follows [1]: ωγ Iγ ¼K ωα Iα
ð1Þ
The macro-hardness was carried out by a HR-150A Rockwell hardness tester with a test load of 150 kg. Ten test points were randomly selected on the surface of the sample, and the average value was calculated as the hardness result. According to ASTM E384-1999 standard, the unnotched impact specimens were tested by a JBW-300 impact tester with 150-J energy at 25 ± 2 °C. The size of samples was 10 mm × 10 mm × 55 mm, and the average of three samples was calculated as the impact toughness result.
Fig. 4. Microstructure of carbidic ductile iron after heat treatment (etched by 4% Nital): (a) OM of CADI after ordinary austempering treatment, (b) SEM image of CADI after ordinary austempering treatment, (c) OM of CADI after S&A treatment, (d) SEM image of CADI after S&A treatment, (e) SEM image at high magnification of CADI after S&A treatment, (f) SEM image at high magnification of CADI after contrast test.
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2.4. Wear test The wear test of CADI was carried out by a MM-200 block-on-ring wear testing machine [14], GCr15 steel with a composition of 0.99C0.35-Mn-0.25Si-1.65Cr- 0.10Mo-0.30Ni-0.25 wt% Cu, a hardness of 60 ± 2 HRC and a dimension of φ40 mm × 10 mm was used as the grinding ring with a rotational speed of 200 r/min and a wear time of 120 min. The size of wear samples was 12.5 mm × 12.5 mm × 19 mm, and the test loads were 100 N, 200 N, 300 N, 400 N, 500 N and 600 N. The weight loss after wear was measured by TG328B balance with 0.01 mg sensitivity, an average weight loss of three samples was calculated as the final result. 3. Results 3.1. Microstructure of as-cast carbidic ductile iron Figs. 2 and 3 show the microstructures and the XRD pattern of ascast carbidic ductile iron with a composition of 3.72C-2.77Si-0.51Mn0.99 wt% Cr respectively. The results indicate that carbidic ductile iron is mainly composed of graphite nodules, eutectic carbides and matrix. Among them, the matrix is lamellar pearlite composed of ferrite and cementite, and the interlaminar spacing of pearlite is about 300 ± 50 nm. According to the results of XRD, the volume fraction and carbon content of retained austenite were estimated by the Eqs. (1) and (2). For the CADI obtained by ordinary austempering treatment, the volume fraction and carbon content of retained austenite are 19 vol% and 1.65 wt %, respectively. For the CADI obtained by S&A treatment, the volume fraction and carbon content of retained austenite are 20.5 vol% and 1.62 wt%, respectively. According to the method in reference [24], the roundness and the percent nodularity of graphite nodules in the microstructure were calculated. The volume fraction of eutectic carbides was measured by Image-pro software. The results show that the percent nodularity of graphite nodules is 92.3%, and the volume fractions of eutectic carbides and graphite in as cast carbidic ductile iron are 13.4 vol%
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and 5.4 vol%, respectively. Due to the addition of Cr into ductile iron, metastable eutectic transformation is promoted, resulting in the formation of eutectic structure containing austenite and eutectic carbides. These eutectic carbides obtained by eutectic reaction are distributed at the grain boundaries and interconnected to form the carbide network, as displayed in Fig. 2b. From the distribution results of EPMA mapping, Cr element is mostly enriched in the carbides (see Fig. 2c). To further determine the type of these eutectic carbides, Fig. 2d illustrates the brightfield TEM micrographs and corresponding selected area diffraction patterns (SADPs) of eutectic carbides. The results indicate that these eutectic carbides are M3C type carbides with orthorhombic lattice belonging to Pnma space group, which has the same lattice type as cementite (Fe3C). However, because a small amount of Cr replaced Fe element in the crystal, these Cr-rich eutectic carbides are called (Fe,Cr)3C [25]. Their lattice constants are a = 0.550153 nm, b = 0.68416 nm, and c = 0.43817 nm, respectively. 3.2. Microstructure of carbidic ductile iron after heat treatment Fig. 4a and b show the optical microstructure and SEM image of carbidic ductile iron after ordinary austempering treatment. Compared with as-cast microstructure, the eutectic carbides changed from network structure to dots, blocks or long shapes, and the corresponding amount were reduced. The matrix changed from lamellar pearlite to ausferrite containing the bainitic ferrite and the retained austenite. From the analysis of image-Pro software, the volume fraction of undissolved eutectic carbides in ordinary CADI is 10 vol%, and the length of ausferrite is about 5–8 μm. Fig. 4c and d show the optical microstructure and SEM image of CADI after S&A treatment (Fig. 1c). The volume fraction of eutectic carbides in CADI obtained by S&A treatment is 2.1 vol%, and the ausferrite is significantly refined (about 1–3 μm), as shown in Fig. 4d. Compared with the XRD results of as-cast carbidic ductile iron (Fig. 3), the austenite phase appeared in the cast iron after S&A treatment. Compared with the ordinary CADI, the volume fraction of eutectic carbides decreased significantly. Fig. 4e shows the SEM image at high
Fig. 5. SEM image after water quenching: (a) SEM image after 1100 °C/10 min + AQ to room temperature + 880 °C/120 min + WQ, (b) SEM image at high magnification after 1100 °C/ 10 min + AQ to room temperature + 880 °C/120 min + WQ, (c) SEM image after 1100 °C/10 min + FC to 880 °C/120 min + WQ (contrast WQ test), (d) SEM image at high magnification after 1100 °C/10 min + FC to 880 °C/120 min + WQ.
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Fig. 6. Bright-field TEM micrographs and the corresponding SADPS of CADI after heat treatment: (a) Bright-field TEM micrographs of CADI after ordinary austempering treatment and corresponding SADPs from bainitic ferrite; (b) Bright-field TEM micrographs of CADI after S&A treatment and corresponding SADPs from precipitate.
magnification of CADI after S&A treatment. There are two types of carbides in the matrix, the undissolved eutectic carbides under high temperature and the precipitated particles after reheating. Interestingly, undissolved eutectic carbides were observed in the matrix after the heat treatment process from Fig. 1b (contrast test), without any precipitated particles, as shown in Fig. 4f. Where, the volume fraction of eutectic carbides in microstructure obtained by contrast test is about 2 vol%. To clearly compare the type and the quantity of two kinds of carbides, the samples were treated with the water quenching (WQ) process from Fig. 1b and c (dashed line), respectively, and the SEM images were shown in Fig. 5. A large amount of precipitated particles appeared in matrix after super-high temperature pretreatment + air quenched (AQ) to room temperature + reheated to 880 °C + water quenched, The total volume fraction of precipitated particles and undissolved eutectic carbides is about 13.3 vol%, and the precipitated particles are mostly distributed around the prior austenite grain boundaries and the undissolved eutectic carbides, as shown in Fig. 5a and b. After superhigh temperature pretreatment + furnace cooled (FC) to 880 °C + water quenched (contrast WQ test), the quantity of eutectic carbides significantly decreased. However, it is noting that no
precipitated particles appeared in matrix after contrast WQ test, as shown in Fig. 5c and d. The CADI obtained by ordinary heat treatment and S&A treatment was compared by a transmission electron microscope, as shown in Fig. 6. For ordinary CADI (Fig. 6a), the white lath is bainitic ferrite, and their thickness is about 100 ± 10 nm. Besides, the retained austenite is distributed between these lath. SADP results show that the bainitic ferrite is cubic lattice (Fig. 6a). Because Si element inhibited the precipitation of ε-carbides, no cementite was observed inside the ferrite. So this bainite was called carbide-free bainite [26–28]. After S&A treatment, the thickness of bainitic ferrite significantly decreased, which was about 50 ± 10 nm, as shown in Fig. 6b. Combined with the EPMA and TEM selected area diffraction pattern (SADP), the chemical compositions and lattice types of precipitated particles were studied. The results of SADP showed that the lattice type of precipitated particles was an orthorhombic lattice, and Cr element was not detected by WDS and EDX in the precipitated particles. In addition, the results of WDS showed that the atomic ratio of Fe to C in the precipitates was about 3:1, indicating that the precipitated particles might be Fe3C with an orthorhombic lattice. Fig. 7 showed the bright-field TEM micrographs and
Fig. 7. Bright-field TEM micrographs and the corresponding SADPs of reflecting the orientation relationship between precipitate and α-Fe: (a) Bright-field TEM micrographs of CADI after S&A treatment, (b) corresponding SADPs of reflecting the orientation relationship between precipitate and α-Fe.
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Fig. 10. Wear loss of CADI under different loads. Fig. 8. Comparison of mechanical properties ordinary austempering treatment and S&A treatment.
the corresponding SADPs used for reflecting the orientation relationship between precipitate and α-Fe. The results indicated that the orientation relationship between bainitic ferrite and precipitated particles may be: (011)α-Fe//(210)Fe3C, and the orientation relationship of zone axis may be: [111]α-Fe//[020]Fe3C. 3.3. Properties of CADI after S&A treatment The Rockwell hardness and impact toughness of CADI after ordinary heat treatment and S&A treatment were compared, and the results were showed in Fig. 8. The results indicate that the hardness of both CADI was similar, but the impact toughness of CADI rose from 9.2. J/cm2 to 19.5 J/ cm2 after S&A treatment, which is increased by 120%. This obvious increase in impact toughness can be correlated to the reduced probability of crack nucleation due to decreasing eutectic carbides and to the reduced stress concentration as a result of breakage of carbide networks [16]. Therefore, the impact toughness of CADI was obviously improved after S&A treatment. According to the heat treatment process from Fig. 1b (contrast test), the contrast CADI was prepared. However, the hardness of CADI obtained by contrast test was significantly lower than ordinary CADI, which was about 43 ± 2 HRC. Cheng et al. [14] found that the Rockwell hardness of CADI gradually increased with the increasing eutectic carbides. The quantity of eutectic carbides was significantly less in the contrast CADI, and there were no precipitated particles. So the hardness was low. After S&A treatment, a large amount of Fe3C particles were precipitated in matrix of CADI, which may be beneficial to the increase of hardness. Therefore, the hardness did not reduce significantly, which is consistent with the precipitation strengthening theory. In addition, the refinement of bainitic ferrite has
a positive effect on the mechanical properties of CADI, which is consistent with the fine-grain strengthening theory [20]. Fig. 9 shows the fracture surface morphology of CADI after impact test. The fracture surface of ordinary CADI is flat, and the cleavage step and microcrack can be observed, indicating that the fracture of ordinary CADI belongs to typical brittle fracture. Researches [15,16] have shown that the cleavage surface of the CADI fracture was formed because the coarse carbides caused stress concentration, promoting the nucleation and propagation of the crack. Therefore, the impact toughness of CADI was only 9.2 J/cm2 after ordinary heat treatment. A large number of dimples were observed in the CADI fracture after S&A treatment, and the precipitated particles were observed inside the dimples. Refined eutectic carbides is helpful to minimize the stress concentration, nucleation and propagation of cracks. In addition, a large amount of particles were precipitated inside the grains, which reduced the carbon content of matrix, resulting in the increase of matrix impact toughness. Therefore, the impact toughness of CADI obtained by S&A treatment sharply increased. Because the precipitated particles with the diameter of about 200 ± 50 nm were fine, the dislocation bypassed the precipitated particles when the crack propagated to the precipitated phase [29], enhancing the impact toughness. The wear loss of two kinds of CADI under different test load was measured in this work, as shown in Fig. 10. The results show that there is no significant difference between two types CADI in the wear loss under low load (below 300 N). However, the CADI obtained by S&A treatment exhibits excellent wear resistance under high load (300–600 N). Many researches [30,31] have shown that the wear resistance of multiphase materials is affected by multiple factors, such as the hardness, volume fraction, shape, size and distribution of the second phase, and matrix properties. After S&A treatment, there are two types of hard phases in the CADI: the undissolved eutectic carbides (Fe, Cr) 3C at the grain boundaries and the precipitated particles (Fe3C) inside the grain. The hardness of (Fe, Cr)3C is slightly higher than Fe3C.
Fig. 9. Fracture morphology of CADI after impact test: (a) SEM image of surface morphology after ordinary austempering treatment; (b) SEM image of surface morphology after S&A treatment.
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Fig. 11. SEM image of carbidic ductile iron after super-high temperature + air cooling: UC: undissolved carbides.
Therefore, although the amount of undissolved eutectic carbides in ordinary CADI (10 vol%) was slightly smaller, the wear resistance was still good under low load. With the increase of load, coarse eutectic carbides was easy to exfoliate under high load, which reduced the wear resistance of ordinary CADI [32]. In addition, fine matrix structure could also improve the wear resistance of materials, which significantly reduced the wear loss of CADI under high load after S&A treatment, improving the wear resistance. 4. Discussions 4.1. Analysis of matrix refinement mechanism To analyze the refinement mechanism of S&A treatment on CADI matrix, the microstructure after super-high temperature + air-cooling (super-high temperature pretreatment) was observed by SEM, as shown in Fig. 11. The black line showed the grain boundaries. Obviously, the high temperature caused the growth of the prior austenite grains, but also increased the supercooling degree (ΔT) of the pearlite transformation. In other words, the thermodynamic driving force of pearlite nucleation increased with the rise of pretreatment temperature, which enhanced the nucleation rate of pearlite, reducing the pearlite interplanar spacing (100 ± 30 nm). This is consistent with previous studies [33,34]. Compared with the as-cast pearlite interlamellar spacing (300 ± 50 nm), the pearlite was obviously refined after hightemperature pretreatment, which also led to the increase of phase boundaries between ferrite and cementite. Fig. 12 shows the schematic
diagram of matrix refinement mechanism. The carbidic ductile iron was reheated to the austenitizing temperature after the super-high temperature pretreatment, the phase transformation from pearlite to austenite began to occur. With the extended holding time, the cementite in the pearlite was gradually dissolved, and the austenite phase interface moved toward the ferrite phase interface. Finally, the prior austenite grains were divided into several austenite grains, as shown in Fig. 12 (the dashed line is new grain boundaries). From the kinetic point of view, the increase of pearlite phase interface was beneficial to the nucleation of austenite, but the growth rate of austenite was slowed down due to the increase of cementite dispersity. From the view of thermodynamic, the growth of austenite needs to span more interface between cementite and ferrite, resulting in the decrease of growth potential energy. Besides, Cr element is carbide forming element, it could slow the growth velocity of austenite grain [35–37]. With the dissolution of eutectic carbides, the Cr diffused from eutectic carbides to matrix, which refined the austenite grain during reheating. Therefore, many austenite grains with small size were formed after reheating, which increased the number of grain boundaries. Because the bainite formation began with nucleation of bainitic ferrite at austenite grain boundaries, the potential sites for bainite formation increased with the increase of the number of quantity boundaries, reducing the length and the thickness of bainitic ferrite. Finally superfine ausferrite matrix is obtained. 4.2. Analysis of Fe3C precipitate phase To study the precipitation behavior of the precipitated phase, the carbidic ductile iron was heated to 1100 °C and then cooled in air and furnace. No precipitate phase was observed in the matrix. Therefore, in order to observe the precipitated particles, the samples were treated by the water quenching process from the dashed line in Fig. 1c. The SEM images were shown in Fig. 13. At 720 °C, the matrix was still pearlite, as shown in Fig. 13a. With the increase of reheating temperature, the cementite in the pearlite gradually decreased. After super-high temperature pretreatment, the content of alloy elements (such as Mn and Cr) in matrix increased with the dissolution of eutectic carbides. Due to the increase of alloy elements in matrix, the eutectoid point moved upward, which caused a part of the cementite in the pearlite to remain. In other words, the precipitated particles might be retained cementite. Fig. 14 shows the bright-field high-resolution TEM (HR-TEM) image and corresponding Fourier-filtered HR-TEM image of the CADI after S&A treatment. The areas A, B, and C in Fig. 14a were α-Fe, Fe3C precipitates, and interfaces between α-Fe and Fe3C, respectively. Fourier-filtered HRTEM image of area A shows that bainitic ferrite is BCC structure. The interplanar distance of the (011) plane for α-Fe is 0.2022 nm, which is in good agreement with the XRD results. The precipitates of Fe3C were orthorhombic lattices, and the interplanar distance of the (011) plane is 0.2067 nm, as shown in Fig. 14c. Noting that dislocations and other defects were rarely found at the interface between α-Fe and Fe3C precipitation, which indicates that the boundaries between bainitic ferrite and Fe3C precipitated particles are coherent boundaries. As is known to all, the coherent boundaries between the precipitated phase and the matrix have a positive effect on improving the properties of materials [38,39]. 5. Conclusions The impact toughness of a 3.72C-2.77Si-0.51Mn-0.99 wt% Cr carbidic ductile iron was improved by the incorporation of a super-high temperature pretreatment prior to austempering treatment (termed here as S&A treatment) without sacrificing hardness and wear resistance. The main conclusions are obtained as follows:
Fig. 12. Schematic diagram of matrix refinement mechanism: P: pearlite, SP: superfine pearlite, γ: austenite, PC: precipitated carbides, AF: ausferrite.
(1) Super-high temperature pretreatment can significantly reduce Cr-rich eutectic carbides at the grain boundary, and promote
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Fig. 13. SEM image of carbidic ductile iron at various reheating temperatures for 120 min + water cooling: (a) 720 °C (b) 760 °C (c) 800 °C (d) 840 °C.
the precipitation of particles inside the prior austenite grains during reheating. The precipitated particles may be Fe3C with the orthorhombic lattice, and the orientation relationship
between bainitic ferrite and precipitated particles may be: (011 )α-Fe//(210)Fe3C. (2) The length of ausferrite is about 1–3 μm, and the thickness of
Fig. 14. High-resolution transmission electron microscopy image and Fourier-filtered images: (a) HR-TEM image of CADI after S&A treatment; (b) Fourier-filtered image of α-Fe in (a); (c) Fourier-filtered image of Fe3C precipitation in (a); (d) Fourier-filtered image of boundary between α-Fe and Fe3C precipitation.
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bainitic ferrite is about 50 ± 10 nm after S&A treatment. This obvious refinement may be attributed to the significant reduction of pearlite lamellar spacing (100 ± 30 nm) after super-high temperature pretreatment and the retardation effect of Cr element on austenite grain growth during reheating. (3) With the rise of reheating temperature from 720 °C to 880 °C, the pearlite decreases gradually, and the precipitated particles are retained. Besides, the boundaries between bainitic ferrite and precipitated particles are coherent boundaries. (4) After S&A treatment, the impact toughness of CADI is 120% higher than that of conventionally heat treated CADI due to the reduction of eutectic carbides. With the formation of precipitated particles, the CADI exhibits excellent wear resistance under high wear load (above 300 N).
CRediT authorship contribution statement Yang Penghui Conceptualization, Investigation, Formal analysis, Writing - original draft. Fu Hanguang Conceptualization, Formal analysis, Writing-original draft, Funding acquisition. Li GuoLu Resources, Writing-review & editing. Liu Jinhai Investigation, Data curation, Project administration, Funding acquisition. Zhao Xuebo Writing - review & editing, Funding acquisition. 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. Acknowledgement The authors would like to thank the financial support for this work from National Natural Science Foundation of China under grant (51775006) and Natural Science Foundation of Hebei Province (E2016202100). References [1] S.K. Putatunda, Development of austempered ductile cast iron (ADI) with simultaneous high yield strength and fracture toughness by a novel two-step austempering process, Mater. Sci. Eng. A 315 (2001) 70–80. [2] H. Zhang, Y.X. Wu, Q.J. Li, X. Hong, Mechanical properties and rolling-sliding wear performance of dual phase austempered ductile iron as potential metro wheel material, Wear 406 (2018) 156–165. [3] A. Meena, M.E. Mansori, Study of dry and minimum quantity lubrication drilling of novel austempered ductile iron (ADI) for automotive applications, Wear 12 (2010) 2412–2416. [4] A. Meena, M.E. Mansori, Drilling performance of green austempered ductile iron (ADI) grade produced by novel manufacturing technology, Int. J. Adv. Manuf. Technol. 59 (2012) 9–19. [5] S. Laino, J.A. Sikora, R.C. Dommarco, Development of wear resistant carbidic austempered ductile iron (CADI), Wear 265 (2008) 1–7. [6] M. Lagarde, A. Basso, J.A. Sikora, R.C. Dommarco, Development and characterization of a new type of ductile iron with a novel multiphase microstructure, ISIJ Int. 51 (2011) 645–650. [7] S. Laino, J.A. Sikora, R.C. Dommarco, Advances in the development of carbidic ADI, Key Eng. Mater. 457 (2011) 187–192. [8] A.S.O. Pimentel, W.L. Guesser, Tratamento térmico de austêmpera em ferro fundido nodular com adições de nióbio e de cromo, Matéria 2 (2017) 22–28. [9] A. Refaey, N. Fatahalla, Effect of microstructure on properties of ADI and low alloyed ductile iron, J. Mater. Sci. 38 (2003) 351–362. [10] C.F. Han, Y.F. Sun, Y. Wu, Y.H. Ma, Effects of vanadium and austempering temperature on microstructure and properties of CADI, Metallogr. Microstruct. Anal. (3) (2015) 135–145.
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Materials and Design journal homepage: www.elsevier.com/locate/matdes
Microstructures and properties of carbidic austempered ductile Iron containing Fe3C particles and superfine ausferrite Yang Penghui a, Fu Hanguang a,⁎, Li Guolu b, Liu Jinhai b, Zhao Xuebo b a b
Key Laboratory of Advanced Functional Materials, Ministry of Education, School of Materials Science and Engineering, Beijing University of Technology, Beijing 100124, PR China School of Materials Science and Engineering, Hebei University of Technology, Tianjin 300130, PR China
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• The impact toughness of CADI is significantly improved by super-high temperature pretreatment and austempering treatment. • The CADI obtained by S&A treatment has excellent wear resistance under high wear load. • With the rise of reheating temperature, the pearlite decreases gradually, and the precipitated particles are retained. • A superfine ausferrite matrix with a length of 1-3 μm and a thickness of 50 ± 10 nm is obtained. • The boundaries between bainitic ferrite and precipitated particles are coherent boundaries.
a r t i c l e
i n f o
Article history: Received 5 June 2019 Received in revised form 22 October 2019 Accepted 13 November 2019 Available online 14 November 2019 Keywords: CADI Super-high temperature pretreatment Superfine ausferrite Precipitates Impact toughness
a b s t r a c t A new type heat treatment process comprising super-high temperature pretreatment and austempering treatment (S&A treatment) was used to process carbidic austempered ductile iron (CADI). The results showed that the netlike eutectic carbides were significantly reduced after super-high temperature pretreatment. Due to the increase in supercooling during super-high temperature pretreatment cooling, the inter-laminar spacing of pearlite was reduced from 300 nm to 100 nm. When the ductile iron with superfine pearlite was reheated, a large amount of Fe3C particles were retained inside the prior austenite grains. After austempering, a superfine ausferrite matrix with a length of 1–3 μm and a thickness of 50 nm was obtained. The main reason for this refinement is the increased number of nucleation-sites for the austenite grains and the retardation-effect of Cr on their growth. TEM results indicated that a possible orientation relationship between bainitic ferrite with a cubic lattice and precipitated Fe3C particles with a orthorhombic lattice is (011)α-Fe//(210)Fe3C, and their boundaries are coherent boundaries. The CADI obtained by S&A treatment has an impact toughness 120% higher than the traditional CADI without sacrificing hardness, and has excellent wear resistance under high wear load. © 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
⁎ Corresponding author at: School of Materials Science and Engineering, Beijing University of Technology, Number 100, Pingle Garden, Chaoyang District, Beijing 100124, PR China. E-mail address: [email protected] (F. Hanguang).
1. Introduction Fracture, corrosion and wear are important failure modes of mechanical parts, which will cause huge economic losses. Therefore, the
https://doi.org/10.1016/j.matdes.2019.108363 0264-1275/© 2019 The Authors. Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
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research on improving the wear resistance and corrosion resistance of materials has been highly valued by researchers. Austempered ductile iron (ADI) is widely used in motor-dom, mining, railway, agricultural machinery and other fields due to its high tensile strength (≤ 1600 MPa), good impact toughness (≥ 100 J/cm2), excellent fatigue resistance and corrosion resistance [1–4]. However, the wear resistance of ADI under high load is not ideal. To improve the wear resistance of ADI, researchers [5–8] invented a new wear-resistant and corrosion-resistant material called Carbidic Austempered Ductile Iron (CADI) on the basis of ADI. There were three main preparation processes for producing CADI: (1) Adding a small amount of alloying elements (such as Nb, V, Mo, Cr, etc.) into ADI to promote the precipitation of eutectic carbides [9–12]; (2) Setting chills in the mould to increase the cooling rate [13]; (3) Reducing the amount of graphitizing elements (in particular Si). Among them, the addition of alloying elements could shorten the crystallization temperature interval between the stable and the metastable eutectic reactions significantly, which promoted the non-equilibrium solidification and facilitated the formation of eutectic carbides. In addition, the addition of Cr could bring more carbides at lower production cost comparing with the same amount of other alloying elements [14]. However, the Cr contained eutectic carbides distributed at grain boundaries, which seriously destroys the continuity of the austenite matrix, resulting in the decrease of impact toughness [14]. Although the effects of alloying elements and cooling rate on hardness, impact toughness, tensile strength and microstructures of CADI have been reported in the past decades [9–14], the important mechanism for improving the toughness of CADI had not been provided. Until 2013, Sun et al. [15,16] proposed a method to improve the impact toughness of CADI using nano-oxides. As a result, when 0.06 wt% nanoceria was added into CADI containing 3.5C-2.85Si-0.75 wt% Cr, the impact toughness was increased by 33% due to the refinement of netlike eutectic carbides. Besides, the abrasive wear rate was reduced by 29%. In addition, the morphology and the size of eutectic carbides were changed by this method, which ignored the refinement of austenite matrix. It has been proved that fine bainitic plates have excellent comprehensive mechanical properties [17–21], such as nanostructured bainitic steel. These bainites are composed of very fine bainitic ferrite plates (about 50 nm) and retained austenite films. Due to the uniform microstructure and the low residual stresses, this bainitic steel has unique advantages in the application of large mechanical components. However, considering the phase transformation kinetics of bainite, the formation of these bainites usually requires long holding time (sometimes longer than several days) at the low temperature, or multi-pass rolling before bainitic transformation [22], which significantly reduces production efficiency. Liu et al. [23] approached the feasibility of adopting super-high austenitizing temperature (960–1000 °C) to improve the impact toughness of CADI. When adding 0.5 wt% Cr and adopting superhigh austenitizing temperature of 960–1000 °C, the impact toughness values (absorbed energy) of CADI specimens with the hardness of 39–50 HRC could reach 28–93 J/cm 2 . However, the grains became coarsened with the increase of austenitizing temperature, and the precipitates had not been found in CADI obtained by super-high austenitizing temperature. When adding 1.0 wt% Cr, the hardness of CADI had exceeded 50 HRC, and the toughness might be under 10 J/cm 2 . Therefore, a method that could greatly improve the impact toughness of CADI without sacrificing wear resistance had not been proposed to date. The present work aims to refine the eutectic carbides and the ausferrite matrix by an innovative heat treatment process including super-high temperature (closed to the melting point) pretreatment and austempering treatment (referred as S&A treatment). The proposed treatment process successfully improves the impact toughness and the wear resistance of CADI owing to the
replacement of the coarse eutectic carbides and coarse ausferrite matrix with lots of precipitated particles and superfine ausferrite, particularly, the wear resistance under medium and high wear load, and the research results are expected to accelerate the popularization and application of CADI, such as digger teeth, balls of ball mill, liners of ball mill, loose soil plow shovels, crusher and some automobile parts et al. 2. Experimental procedures 2.1. Material preparation and heat treatment The charge materials were melted using a 50 kg medium frequency induction furnace. 0.15 wt% pure aluminum was added to the melt to deoxidize the molten iron at 1480–1500 °C, and then transferred to a ladle containing 1.2 wt% inoculant and 1.5 wt% nodularizer. Spheroidization and inoculation were performed by pour-over treatment process, and inoculant and nodularizer were FeSi75 (75 wt% Si) and FeSiMg6Re2 (6Mg-2Re-40 wt% Si), respectively. After removing the slag, the melt was poured into sodium silicate-CO2 bonded sand molds, obtaining Y-block ingots according to ASTM A781/A 781-M95. The chemical composition of the ingots was measured by QualiSpark 750 Optical Emission Spectrometer, which was 3.72C-2.77Si-0.51Mn-0.99 wt% Cr. Samples with dimensions of 11 mm × 11 mm × 55 mm and 13 mm × 13 mm × 19 mm were cut from the bottom of Y- block ingots for impact test, microstructural analysis and wear test. Fig. 1a shows a traditional austempering treatment. The operation processes are as follows: The samples were heated to 880 °C for 120 min, and then instantly quenched into salt-bath of 50 wt% KNO350 wt% NaNO3 and isothermally held for 120 min at 300 °C, followed by air-cooling, called “Ordinary austempering treatment” below. Fig. 1b shows a super-high temperature treatment, Here, the sample was heated to 1100 °C for 10 min, cooled to 880 °C in the furnace, and held for 120 min, followed by isothermal quenching, called “Contrast test”. Fig. 1b also shows a water quenching treatment (dashed line), and the operation processes are as follows: The sample was heated to 1100 °C for 10 min, cooled to 880 °C in the furnace, and held for 120 min before water-quenched, called “Contrast WQ test”. To improve the mechanical properties of CADI and refine the ausferrite, a heat treatment process including super-high temperature pretreatment and austempering treatment used in the present work is shown in Fig. 1c (black lines), called “S&A treatment” below. The specific operation processes are as follows: Firstly, the sample was heated to 1100 °C for 10 min, then air cooled to room temperature. Secondly, the samples were reheated to 880 °C for 120 min, and then quenched into saltbath. Finally, isothermally held for 120 min at 300 °C, followed by aircooling. To more conveniently observe the precipitated particles, these samples were reheated to 720 °C, 760 °C, 800 °C, 840 °C, and 880 °C respectively after super-high temperature pretreatment with a duration of 120 min, followed by water-quenched. 2.2. Microstructural analysis The microstructures were investigated by an Olympus BX51 optical microscope (OM), a JSM-6510 scanning electron microscope (SEM), a Rigaku-Utima-IV X-ray diffraction (XRD), a JEM-2100F transmission electron microscopy (TEM) with an energy dispersive X-ray (EDX), an electron probe micro-analyzer (EPMA) with a wavelength dispersive X-ray (WDS). Cu-Kα was used to perform the radiation at 40 kV and 30 mA. The specimens were scanned in the angular 2θ range from 20° to 90° with a step size of 0.5°/min and a collection time of 10 s. TEM samples preparation processes were as follows: firstly, the samples with an area of 100 mm 2 and a thickness of 0.5 mm were cut by wire cutting; then mechanically ground by hand to a thickness of 30–50 μm without cracks and
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Fig. 1. Schematic diagram of heat treatments; (a) ordinary austempering treatment, (b) contrast test, (c) S&A treatment. SQ: salt quenching, AQ: air quenching, FC: furnace cooling, WQ: water quenching.
breakage; finally, thinning by plasma at 4.5 kV, 1.9 mA and rotating angle of 6°. The area of precipitated particles and eutectic carbides was calculated by Image-pro software. 20 fields of view were
randomly chosen, and the average value was calculated as the final result. The volume fraction and carbon content of retained austenite were estimated by XRD. The calculation formula of
Fig. 2. Microstructure of as-cast carbidic ductile iron: (a) optical microstructure (etched by 4% Nital), (b) SEM image (etched by 4% Nital), (c) Mapping distribution of Cr elements, (d) Bright-field TEM micrographs and corresponding selected area diffraction patterns (SADPs) of eutectic carbides.
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where ω γ is the volume fractions of austenite, ω α is the volume fractions of ferrite, I γ is the integral intensities of austenite, I α is the integral intensities of ferrite. K = 1.39. The carbon content of retained austenite is calculated by the following Eq. (2) [2]: aγ ¼ 0:3548 þ 0:0044 C γ
ð2Þ
where Cγ is the carbon content of retained austenite (wt%), aγ is the lattice constants of austenite (Å). 2.3. Mechanical performance test Fig. 3. XRD patterns of as-cast carbidic ductile iron, ordinary CADI and CADI of after S&A treatment: (a) as-cast, (b) ordinary CADI, (c) S&A treatment.
volume fraction is as follows [1]: ωγ Iγ ¼K ωα Iα
ð1Þ
The macro-hardness was carried out by a HR-150A Rockwell hardness tester with a test load of 150 kg. Ten test points were randomly selected on the surface of the sample, and the average value was calculated as the hardness result. According to ASTM E384-1999 standard, the unnotched impact specimens were tested by a JBW-300 impact tester with 150-J energy at 25 ± 2 °C. The size of samples was 10 mm × 10 mm × 55 mm, and the average of three samples was calculated as the impact toughness result.
Fig. 4. Microstructure of carbidic ductile iron after heat treatment (etched by 4% Nital): (a) OM of CADI after ordinary austempering treatment, (b) SEM image of CADI after ordinary austempering treatment, (c) OM of CADI after S&A treatment, (d) SEM image of CADI after S&A treatment, (e) SEM image at high magnification of CADI after S&A treatment, (f) SEM image at high magnification of CADI after contrast test.
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2.4. Wear test The wear test of CADI was carried out by a MM-200 block-on-ring wear testing machine [14], GCr15 steel with a composition of 0.99C0.35-Mn-0.25Si-1.65Cr- 0.10Mo-0.30Ni-0.25 wt% Cu, a hardness of 60 ± 2 HRC and a dimension of φ40 mm × 10 mm was used as the grinding ring with a rotational speed of 200 r/min and a wear time of 120 min. The size of wear samples was 12.5 mm × 12.5 mm × 19 mm, and the test loads were 100 N, 200 N, 300 N, 400 N, 500 N and 600 N. The weight loss after wear was measured by TG328B balance with 0.01 mg sensitivity, an average weight loss of three samples was calculated as the final result. 3. Results 3.1. Microstructure of as-cast carbidic ductile iron Figs. 2 and 3 show the microstructures and the XRD pattern of ascast carbidic ductile iron with a composition of 3.72C-2.77Si-0.51Mn0.99 wt% Cr respectively. The results indicate that carbidic ductile iron is mainly composed of graphite nodules, eutectic carbides and matrix. Among them, the matrix is lamellar pearlite composed of ferrite and cementite, and the interlaminar spacing of pearlite is about 300 ± 50 nm. According to the results of XRD, the volume fraction and carbon content of retained austenite were estimated by the Eqs. (1) and (2). For the CADI obtained by ordinary austempering treatment, the volume fraction and carbon content of retained austenite are 19 vol% and 1.65 wt %, respectively. For the CADI obtained by S&A treatment, the volume fraction and carbon content of retained austenite are 20.5 vol% and 1.62 wt%, respectively. According to the method in reference [24], the roundness and the percent nodularity of graphite nodules in the microstructure were calculated. The volume fraction of eutectic carbides was measured by Image-pro software. The results show that the percent nodularity of graphite nodules is 92.3%, and the volume fractions of eutectic carbides and graphite in as cast carbidic ductile iron are 13.4 vol%
5
and 5.4 vol%, respectively. Due to the addition of Cr into ductile iron, metastable eutectic transformation is promoted, resulting in the formation of eutectic structure containing austenite and eutectic carbides. These eutectic carbides obtained by eutectic reaction are distributed at the grain boundaries and interconnected to form the carbide network, as displayed in Fig. 2b. From the distribution results of EPMA mapping, Cr element is mostly enriched in the carbides (see Fig. 2c). To further determine the type of these eutectic carbides, Fig. 2d illustrates the brightfield TEM micrographs and corresponding selected area diffraction patterns (SADPs) of eutectic carbides. The results indicate that these eutectic carbides are M3C type carbides with orthorhombic lattice belonging to Pnma space group, which has the same lattice type as cementite (Fe3C). However, because a small amount of Cr replaced Fe element in the crystal, these Cr-rich eutectic carbides are called (Fe,Cr)3C [25]. Their lattice constants are a = 0.550153 nm, b = 0.68416 nm, and c = 0.43817 nm, respectively. 3.2. Microstructure of carbidic ductile iron after heat treatment Fig. 4a and b show the optical microstructure and SEM image of carbidic ductile iron after ordinary austempering treatment. Compared with as-cast microstructure, the eutectic carbides changed from network structure to dots, blocks or long shapes, and the corresponding amount were reduced. The matrix changed from lamellar pearlite to ausferrite containing the bainitic ferrite and the retained austenite. From the analysis of image-Pro software, the volume fraction of undissolved eutectic carbides in ordinary CADI is 10 vol%, and the length of ausferrite is about 5–8 μm. Fig. 4c and d show the optical microstructure and SEM image of CADI after S&A treatment (Fig. 1c). The volume fraction of eutectic carbides in CADI obtained by S&A treatment is 2.1 vol%, and the ausferrite is significantly refined (about 1–3 μm), as shown in Fig. 4d. Compared with the XRD results of as-cast carbidic ductile iron (Fig. 3), the austenite phase appeared in the cast iron after S&A treatment. Compared with the ordinary CADI, the volume fraction of eutectic carbides decreased significantly. Fig. 4e shows the SEM image at high
Fig. 5. SEM image after water quenching: (a) SEM image after 1100 °C/10 min + AQ to room temperature + 880 °C/120 min + WQ, (b) SEM image at high magnification after 1100 °C/ 10 min + AQ to room temperature + 880 °C/120 min + WQ, (c) SEM image after 1100 °C/10 min + FC to 880 °C/120 min + WQ (contrast WQ test), (d) SEM image at high magnification after 1100 °C/10 min + FC to 880 °C/120 min + WQ.
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Fig. 6. Bright-field TEM micrographs and the corresponding SADPS of CADI after heat treatment: (a) Bright-field TEM micrographs of CADI after ordinary austempering treatment and corresponding SADPs from bainitic ferrite; (b) Bright-field TEM micrographs of CADI after S&A treatment and corresponding SADPs from precipitate.
magnification of CADI after S&A treatment. There are two types of carbides in the matrix, the undissolved eutectic carbides under high temperature and the precipitated particles after reheating. Interestingly, undissolved eutectic carbides were observed in the matrix after the heat treatment process from Fig. 1b (contrast test), without any precipitated particles, as shown in Fig. 4f. Where, the volume fraction of eutectic carbides in microstructure obtained by contrast test is about 2 vol%. To clearly compare the type and the quantity of two kinds of carbides, the samples were treated with the water quenching (WQ) process from Fig. 1b and c (dashed line), respectively, and the SEM images were shown in Fig. 5. A large amount of precipitated particles appeared in matrix after super-high temperature pretreatment + air quenched (AQ) to room temperature + reheated to 880 °C + water quenched, The total volume fraction of precipitated particles and undissolved eutectic carbides is about 13.3 vol%, and the precipitated particles are mostly distributed around the prior austenite grain boundaries and the undissolved eutectic carbides, as shown in Fig. 5a and b. After superhigh temperature pretreatment + furnace cooled (FC) to 880 °C + water quenched (contrast WQ test), the quantity of eutectic carbides significantly decreased. However, it is noting that no
precipitated particles appeared in matrix after contrast WQ test, as shown in Fig. 5c and d. The CADI obtained by ordinary heat treatment and S&A treatment was compared by a transmission electron microscope, as shown in Fig. 6. For ordinary CADI (Fig. 6a), the white lath is bainitic ferrite, and their thickness is about 100 ± 10 nm. Besides, the retained austenite is distributed between these lath. SADP results show that the bainitic ferrite is cubic lattice (Fig. 6a). Because Si element inhibited the precipitation of ε-carbides, no cementite was observed inside the ferrite. So this bainite was called carbide-free bainite [26–28]. After S&A treatment, the thickness of bainitic ferrite significantly decreased, which was about 50 ± 10 nm, as shown in Fig. 6b. Combined with the EPMA and TEM selected area diffraction pattern (SADP), the chemical compositions and lattice types of precipitated particles were studied. The results of SADP showed that the lattice type of precipitated particles was an orthorhombic lattice, and Cr element was not detected by WDS and EDX in the precipitated particles. In addition, the results of WDS showed that the atomic ratio of Fe to C in the precipitates was about 3:1, indicating that the precipitated particles might be Fe3C with an orthorhombic lattice. Fig. 7 showed the bright-field TEM micrographs and
Fig. 7. Bright-field TEM micrographs and the corresponding SADPs of reflecting the orientation relationship between precipitate and α-Fe: (a) Bright-field TEM micrographs of CADI after S&A treatment, (b) corresponding SADPs of reflecting the orientation relationship between precipitate and α-Fe.
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Fig. 10. Wear loss of CADI under different loads. Fig. 8. Comparison of mechanical properties ordinary austempering treatment and S&A treatment.
the corresponding SADPs used for reflecting the orientation relationship between precipitate and α-Fe. The results indicated that the orientation relationship between bainitic ferrite and precipitated particles may be: (011)α-Fe//(210)Fe3C, and the orientation relationship of zone axis may be: [111]α-Fe//[020]Fe3C. 3.3. Properties of CADI after S&A treatment The Rockwell hardness and impact toughness of CADI after ordinary heat treatment and S&A treatment were compared, and the results were showed in Fig. 8. The results indicate that the hardness of both CADI was similar, but the impact toughness of CADI rose from 9.2. J/cm2 to 19.5 J/ cm2 after S&A treatment, which is increased by 120%. This obvious increase in impact toughness can be correlated to the reduced probability of crack nucleation due to decreasing eutectic carbides and to the reduced stress concentration as a result of breakage of carbide networks [16]. Therefore, the impact toughness of CADI was obviously improved after S&A treatment. According to the heat treatment process from Fig. 1b (contrast test), the contrast CADI was prepared. However, the hardness of CADI obtained by contrast test was significantly lower than ordinary CADI, which was about 43 ± 2 HRC. Cheng et al. [14] found that the Rockwell hardness of CADI gradually increased with the increasing eutectic carbides. The quantity of eutectic carbides was significantly less in the contrast CADI, and there were no precipitated particles. So the hardness was low. After S&A treatment, a large amount of Fe3C particles were precipitated in matrix of CADI, which may be beneficial to the increase of hardness. Therefore, the hardness did not reduce significantly, which is consistent with the precipitation strengthening theory. In addition, the refinement of bainitic ferrite has
a positive effect on the mechanical properties of CADI, which is consistent with the fine-grain strengthening theory [20]. Fig. 9 shows the fracture surface morphology of CADI after impact test. The fracture surface of ordinary CADI is flat, and the cleavage step and microcrack can be observed, indicating that the fracture of ordinary CADI belongs to typical brittle fracture. Researches [15,16] have shown that the cleavage surface of the CADI fracture was formed because the coarse carbides caused stress concentration, promoting the nucleation and propagation of the crack. Therefore, the impact toughness of CADI was only 9.2 J/cm2 after ordinary heat treatment. A large number of dimples were observed in the CADI fracture after S&A treatment, and the precipitated particles were observed inside the dimples. Refined eutectic carbides is helpful to minimize the stress concentration, nucleation and propagation of cracks. In addition, a large amount of particles were precipitated inside the grains, which reduced the carbon content of matrix, resulting in the increase of matrix impact toughness. Therefore, the impact toughness of CADI obtained by S&A treatment sharply increased. Because the precipitated particles with the diameter of about 200 ± 50 nm were fine, the dislocation bypassed the precipitated particles when the crack propagated to the precipitated phase [29], enhancing the impact toughness. The wear loss of two kinds of CADI under different test load was measured in this work, as shown in Fig. 10. The results show that there is no significant difference between two types CADI in the wear loss under low load (below 300 N). However, the CADI obtained by S&A treatment exhibits excellent wear resistance under high load (300–600 N). Many researches [30,31] have shown that the wear resistance of multiphase materials is affected by multiple factors, such as the hardness, volume fraction, shape, size and distribution of the second phase, and matrix properties. After S&A treatment, there are two types of hard phases in the CADI: the undissolved eutectic carbides (Fe, Cr) 3C at the grain boundaries and the precipitated particles (Fe3C) inside the grain. The hardness of (Fe, Cr)3C is slightly higher than Fe3C.
Fig. 9. Fracture morphology of CADI after impact test: (a) SEM image of surface morphology after ordinary austempering treatment; (b) SEM image of surface morphology after S&A treatment.
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Fig. 11. SEM image of carbidic ductile iron after super-high temperature + air cooling: UC: undissolved carbides.
Therefore, although the amount of undissolved eutectic carbides in ordinary CADI (10 vol%) was slightly smaller, the wear resistance was still good under low load. With the increase of load, coarse eutectic carbides was easy to exfoliate under high load, which reduced the wear resistance of ordinary CADI [32]. In addition, fine matrix structure could also improve the wear resistance of materials, which significantly reduced the wear loss of CADI under high load after S&A treatment, improving the wear resistance. 4. Discussions 4.1. Analysis of matrix refinement mechanism To analyze the refinement mechanism of S&A treatment on CADI matrix, the microstructure after super-high temperature + air-cooling (super-high temperature pretreatment) was observed by SEM, as shown in Fig. 11. The black line showed the grain boundaries. Obviously, the high temperature caused the growth of the prior austenite grains, but also increased the supercooling degree (ΔT) of the pearlite transformation. In other words, the thermodynamic driving force of pearlite nucleation increased with the rise of pretreatment temperature, which enhanced the nucleation rate of pearlite, reducing the pearlite interplanar spacing (100 ± 30 nm). This is consistent with previous studies [33,34]. Compared with the as-cast pearlite interlamellar spacing (300 ± 50 nm), the pearlite was obviously refined after hightemperature pretreatment, which also led to the increase of phase boundaries between ferrite and cementite. Fig. 12 shows the schematic
diagram of matrix refinement mechanism. The carbidic ductile iron was reheated to the austenitizing temperature after the super-high temperature pretreatment, the phase transformation from pearlite to austenite began to occur. With the extended holding time, the cementite in the pearlite was gradually dissolved, and the austenite phase interface moved toward the ferrite phase interface. Finally, the prior austenite grains were divided into several austenite grains, as shown in Fig. 12 (the dashed line is new grain boundaries). From the kinetic point of view, the increase of pearlite phase interface was beneficial to the nucleation of austenite, but the growth rate of austenite was slowed down due to the increase of cementite dispersity. From the view of thermodynamic, the growth of austenite needs to span more interface between cementite and ferrite, resulting in the decrease of growth potential energy. Besides, Cr element is carbide forming element, it could slow the growth velocity of austenite grain [35–37]. With the dissolution of eutectic carbides, the Cr diffused from eutectic carbides to matrix, which refined the austenite grain during reheating. Therefore, many austenite grains with small size were formed after reheating, which increased the number of grain boundaries. Because the bainite formation began with nucleation of bainitic ferrite at austenite grain boundaries, the potential sites for bainite formation increased with the increase of the number of quantity boundaries, reducing the length and the thickness of bainitic ferrite. Finally superfine ausferrite matrix is obtained. 4.2. Analysis of Fe3C precipitate phase To study the precipitation behavior of the precipitated phase, the carbidic ductile iron was heated to 1100 °C and then cooled in air and furnace. No precipitate phase was observed in the matrix. Therefore, in order to observe the precipitated particles, the samples were treated by the water quenching process from the dashed line in Fig. 1c. The SEM images were shown in Fig. 13. At 720 °C, the matrix was still pearlite, as shown in Fig. 13a. With the increase of reheating temperature, the cementite in the pearlite gradually decreased. After super-high temperature pretreatment, the content of alloy elements (such as Mn and Cr) in matrix increased with the dissolution of eutectic carbides. Due to the increase of alloy elements in matrix, the eutectoid point moved upward, which caused a part of the cementite in the pearlite to remain. In other words, the precipitated particles might be retained cementite. Fig. 14 shows the bright-field high-resolution TEM (HR-TEM) image and corresponding Fourier-filtered HR-TEM image of the CADI after S&A treatment. The areas A, B, and C in Fig. 14a were α-Fe, Fe3C precipitates, and interfaces between α-Fe and Fe3C, respectively. Fourier-filtered HRTEM image of area A shows that bainitic ferrite is BCC structure. The interplanar distance of the (011) plane for α-Fe is 0.2022 nm, which is in good agreement with the XRD results. The precipitates of Fe3C were orthorhombic lattices, and the interplanar distance of the (011) plane is 0.2067 nm, as shown in Fig. 14c. Noting that dislocations and other defects were rarely found at the interface between α-Fe and Fe3C precipitation, which indicates that the boundaries between bainitic ferrite and Fe3C precipitated particles are coherent boundaries. As is known to all, the coherent boundaries between the precipitated phase and the matrix have a positive effect on improving the properties of materials [38,39]. 5. Conclusions The impact toughness of a 3.72C-2.77Si-0.51Mn-0.99 wt% Cr carbidic ductile iron was improved by the incorporation of a super-high temperature pretreatment prior to austempering treatment (termed here as S&A treatment) without sacrificing hardness and wear resistance. The main conclusions are obtained as follows:
Fig. 12. Schematic diagram of matrix refinement mechanism: P: pearlite, SP: superfine pearlite, γ: austenite, PC: precipitated carbides, AF: ausferrite.
(1) Super-high temperature pretreatment can significantly reduce Cr-rich eutectic carbides at the grain boundary, and promote
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Fig. 13. SEM image of carbidic ductile iron at various reheating temperatures for 120 min + water cooling: (a) 720 °C (b) 760 °C (c) 800 °C (d) 840 °C.
the precipitation of particles inside the prior austenite grains during reheating. The precipitated particles may be Fe3C with the orthorhombic lattice, and the orientation relationship
between bainitic ferrite and precipitated particles may be: (011 )α-Fe//(210)Fe3C. (2) The length of ausferrite is about 1–3 μm, and the thickness of
Fig. 14. High-resolution transmission electron microscopy image and Fourier-filtered images: (a) HR-TEM image of CADI after S&A treatment; (b) Fourier-filtered image of α-Fe in (a); (c) Fourier-filtered image of Fe3C precipitation in (a); (d) Fourier-filtered image of boundary between α-Fe and Fe3C precipitation.
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bainitic ferrite is about 50 ± 10 nm after S&A treatment. This obvious refinement may be attributed to the significant reduction of pearlite lamellar spacing (100 ± 30 nm) after super-high temperature pretreatment and the retardation effect of Cr element on austenite grain growth during reheating. (3) With the rise of reheating temperature from 720 °C to 880 °C, the pearlite decreases gradually, and the precipitated particles are retained. Besides, the boundaries between bainitic ferrite and precipitated particles are coherent boundaries. (4) After S&A treatment, the impact toughness of CADI is 120% higher than that of conventionally heat treated CADI due to the reduction of eutectic carbides. With the formation of precipitated particles, the CADI exhibits excellent wear resistance under high wear load (above 300 N).
CRediT authorship contribution statement Yang Penghui Conceptualization, Investigation, Formal analysis, Writing - original draft. Fu Hanguang Conceptualization, Formal analysis, Writing-original draft, Funding acquisition. Li GuoLu Resources, Writing-review & editing. Liu Jinhai Investigation, Data curation, Project administration, Funding acquisition. Zhao Xuebo Writing - review & editing, Funding acquisition. 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. Acknowledgement The authors would like to thank the financial support for this work from National Natural Science Foundation of China under grant (51775006) and Natural Science Foundation of Hebei Province (E2016202100). References [1] S.K. Putatunda, Development of austempered ductile cast iron (ADI) with simultaneous high yield strength and fracture toughness by a novel two-step austempering process, Mater. Sci. Eng. A 315 (2001) 70–80. [2] H. Zhang, Y.X. Wu, Q.J. Li, X. Hong, Mechanical properties and rolling-sliding wear performance of dual phase austempered ductile iron as potential metro wheel material, Wear 406 (2018) 156–165. [3] A. Meena, M.E. Mansori, Study of dry and minimum quantity lubrication drilling of novel austempered ductile iron (ADI) for automotive applications, Wear 12 (2010) 2412–2416. [4] A. Meena, M.E. Mansori, Drilling performance of green austempered ductile iron (ADI) grade produced by novel manufacturing technology, Int. J. Adv. Manuf. Technol. 59 (2012) 9–19. [5] S. Laino, J.A. Sikora, R.C. Dommarco, Development of wear resistant carbidic austempered ductile iron (CADI), Wear 265 (2008) 1–7. [6] M. Lagarde, A. Basso, J.A. Sikora, R.C. Dommarco, Development and characterization of a new type of ductile iron with a novel multiphase microstructure, ISIJ Int. 51 (2011) 645–650. [7] S. Laino, J.A. Sikora, R.C. Dommarco, Advances in the development of carbidic ADI, Key Eng. Mater. 457 (2011) 187–192. [8] A.S.O. Pimentel, W.L. Guesser, Tratamento térmico de austêmpera em ferro fundido nodular com adições de nióbio e de cromo, Matéria 2 (2017) 22–28. [9] A. Refaey, N. Fatahalla, Effect of microstructure on properties of ADI and low alloyed ductile iron, J. Mater. Sci. 38 (2003) 351–362. [10] C.F. Han, Y.F. Sun, Y. Wu, Y.H. Ma, Effects of vanadium and austempering temperature on microstructure and properties of CADI, Metallogr. Microstruct. Anal. (3) (2015) 135–145.
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