Issue
Metall. Res. Technol.
Volume 119, Number 2, 2022
Article Number 211
Number of page(s) 13
DOI https://doi.org/10.1051/metal/2022025
Published online 27 April 2022

© EDP Sciences, 2022

1 Introduction

High-speed steels (HSS) have been extensively developed as cutting tool steels, practically in the form of hobs and broaches in gear cutting machines. Usually, these steels are tempered after marquenching treatment, which results in a good combination of high wear resistance and fracture toughness, especially at high temperatures [1]. These two properties depend, on the one hand, on the chemical compositions and microstructural parameters such as the primary austenite grain size and the size and distribution of the primary coarse carbides formed during the manufacturing process of the raw materials for tool steels in the annealed or hot-worked condition [2], and, on the other hand, on the final hardening heat treatment that determines the final microstructure of the steel. Although a worn tool can be restored by grinding, a cracked tool has to be scrapped, which is costly for the manufacturers. In this context, the study of fracture behavior, especially the fracture toughness of high-speed steels, is of particular importance.

Lee et al. [3] studied the influence of microstructural parameters (e.g. microstructural constituents and the morphology of primary carbide particles) on the fracture properties of high-speed steels. They reported that microcracks initially form at the primary carbides and then propagate rapidly along the intercellular networks. Sarafianos [4] in his study on HSS also found that cracks form at the interface between the matrix and large carbides. The effect of heat treatment parameters such as austenitizing and tempering time and temperature on the fracture toughness of marquenched tool steels has been investigated in several studies. For example, in an early work by Crane and Bigg [5], they studied the fracture toughness of AISIM2, T1 and A2 tool steels using different heat treatment cycles in the marquenched condition required for maximum toughness without loss of hardness. After changing the temperature and duration of the austenitizing and tempering stages, they showed that toughness improved slightly when the hardness was reduced to about 520 HV. In another study by Cajner et al. [6], the effect of deep cryogenic treatment on the fracture toughness of a heat treated powder metallurgical high speed steel was investigated. Although this material exhibited a slight increase in hardness and higher wear resistance due to the formation of η-carbides after tempering, the effect of deep cryogenic treatment on its fracture toughness was not significant. Tu et al [7] showed that for JIS SK5 steel, marquencing can be replaced by the austempering process, which converts the austenite to bainite and replaces the martensitic microstructure to improve the fracture toughness. In a study by Teker and Yilmaz [8] on the effect of austempering on the mechanical properties of solution annealed ferritic nodular cast iron, they have shown that the ausferrite in the microstructure plays a positive role in ductility, toughness and strain rates. Kaleicheva [9] studied the mechanism and kinetics of bainitic transformation in HS18-0-1 high speed steel. He pointed out that the austenitizing temperature has no effect on the mechanism of bainitic transformation, but the transformation kinetics of the austenitizing process increases with increasing austenitizing temperature. In a preliminary study by Xuejie and Tianjian [10], they investigated the correlation of bainitic and martensitic microstructures with secondary hardening after tempering of EN ISO 4957, HS 6-5-4-2 high-speed steel. They found that secondary hardening affects softening resistance along with notched impact strength, which is higher for the bainitic microstructure than for the martensitic structure. These studies are limited to changing the process parameters (i.e. time and temperature during austenitization and tempering) and achieve high toughness in the tempered martensitic state. Güral and Altuntas [11] performed a sequence of heat treatments on PM Fe-C material that included austenitization at 950 °C followed by spherodization at 705 °C. The samples were then partially austenitized at 750 °C and austempered at 350 °C. They concluded that the remaining carbide due to partial dissolution and the binitic microstructure increased the impact strength of the samples compared to the quenched samples.

To the best of the authors’ knowledge, there are no reports on the evaluation of the fracture behavior of these steels in tempered bainitic microstructure conditions. In order to gain insight into the fracture behavior of high-speed steels, a series of marquenching processes were performed on PM HSS ASP2030 and their fracture toughness was compared with that obtained by the austenitizing austempering process at different austenitizing temperatures (1150–1185 °C) with the highest hardness value between 65 and 66 HRC.

In most reports, circumferentially notched tensile (CNT) specimens were used instead of conventional CT specimens to take advantage of the radial symmetry of heat transfer to the CNT specimen, which allows the formation of a completely uniform microstructure [12]. To evaluate the fracture toughness, CNT specimens without pre-crack were used and the microstructural features, including the size and volume fraction of the undissolved carbides and prior austenite grain size, together with the dislocation density and volume fraction of the retained austenite, were investigated using scanning electron microscopy and X-ray diffraction analysis. Since identification of the matrix microstructure was almost impossible even at higher magnifications using the images from Field Emission Scanning Electron Microscopes (FE-SEM), the width of the ferrite laths used as a means of evaluation by Electron Backscatter Diffraction (EBSD) experiments.

2 Experimental procedures

2.1 Material and heat treatment

In this study, a cylindrical rod of raw material PM HSS ASP2030 with a diameter of 20 mm was used. The steel was produced by a powder metallurgical process followed by hot isostatic pressing (HIP). The bar was then cut into various lengths of 20 mm diameter for SEM and fracture testing using a wire electrical discharge machine (WEDM). The chemical composition given in Table 1 was determined using a spark emission spectrometer.

Dilatometry tests were performed to determine the critical temperatures of the steel and to develop appropriate heat treatment cycles. Figure 1 illustrates the results of the dilatometry tests.

AC1 and AC3 temperatures of the primary steel were determined to be 837.4 and 860.5 °C, respectively. Accordingly, the specimens were divided into two groups for marquenching and austempering heat treatments. The heat treatment processes are schematically shown in Figure 2. In order to obtain a uniform temperature distribution on the surface and in the core of all samples in the austenitizing stage, the samples were preheated in two stages. In the first stage, they were held at 550 °C for 30 minutes in a Dequssa™ muffle furnace and then immersed in a molten salt bath at 850 °C for 15 minutes in the second stage. The samples were then austenitized at 1150, 1160, 1170 and 1185 °C for 150 seconds. The austenitizing temperatures range were chozen based on the ErasteelTM [13] recommendation for similar composition high speed steels and to obtain a combination of high hardness and sufficient toughness [5].

In the marquenching process, after austenitizing, the samples were quenched to 510 ± 5 °C and held at this temperature for 8 minutes, followed by air cooling to room temperature at a cooling rate of 1.0 °C/S. In the austempering process, samples first quenched to 510 ± 5 °C similarly by air cooling from the austenitizing temperature and held for 60 min and then austempered in a molten salt bath at 235 °C for 72 h and then cooled in air. A longer time for austempering was chosen according to Güral and Altuntas [11], who reported that this leads to higher fracture toughness due to the formation of a larger volume fraction of bainite.

Tated and Patil [14] have shown that the retained austenite transforms to martensite under working conditions due to transformation induced plasticity, resulting in reduced tool life. This makes the steel more susceptible to dimensional instability in service.

Therefore, to minimize the volume fraction of retained austenite in all samples, triple tempering at 560 °C for 2 h was performed as suggested by Vilar et al. [15].

To denote all the samples in this article, the marquenched and austempered samples were denoted by the symbols M and A, respectively, followed by two numbers separated by a dashed line. The first number indicates the residence time in the quenching medium in minutes and the second number indicates the austenitizing temperature in °C. For example, A60-1170 represents the austenitized sample held at 510 °C for 60 min after austenitization at 1170 °C. The preliminary results of this study showed that increasing the holding time at 510 °C for 60 min had little effect on the fracture behavior of the M60 samples compared to the M8 samples, while it improved the kinetics of bainitic transformation in the tool steel PM HSS ASP2030. Some of these results are presented here, the rest will be reported in a separate article.

Table 1

Chemical composition of PM HSS ASP2030 (wt.%).

thumbnail Fig. 1

Dilatometry curve to measure the AC1 and AC3 temperatures.

thumbnail Fig. 2

Schematic illustration of austempering and marquenching heat treatment cycles.

2.2 Microstructure characterization

The cross section of the specimens perpendicular to the HIP direction was used for microstructural studies. The specimens were mechanically polished and chemically etched with the Vilella’s reagent (1 g picric acid + 5 mL hydrochloric acid + 95 mL ethanol) for 30–60 s to study the microstructures.

Characterization of the microstructure of the samples was performed using a Nova NanoSEM450™ field emission scanning electron microscope (FE-SEM) operated at 20 kV and a working distance of 11.3 mm. Image processing was performed using Clemex Vision PE® software to determine the volume fraction and size of undissolved carbides observed by SEM microscopy after heat treatment cycles. To measure the prior austenite grain size, equivalent circles of the smallest and largest diameters were inscribed and circumscribed, respectively, according to the ASTM E112 standard. A minimum of 10 microscopic images were taken at 3000× magnification and the average value was then reported with an uncertainty of 95%. Samples were prepared using standard metallographic procedures, including grinding and diamond polishing and final polishing with colloidal SiO2 for 12 h. A field emission scanning electron microscope (FEI Quanta FEG 450) with an EBSD detector was used to characterize the microstructure of the tempered samples. The EBSD maps were acquired with a step size of 60 nm in the 50 × 50 μm2 scaned area and processed using EDAX-TSLOIM AnalysisTM software version 7.2.1. The grains were defined by a 5-degree misorientation threshold and a minimum number of 2 pixels. Grain size was measured based on the linear intercept method and by averaging the lines in vertical and horizontal directions [16] according to ASTM E112 standard [17] for bainitic-martensite laths.

The X-ray diffraction patterns were collected to characterize the phases present, i.e. position and broadening of the peaks and also to measure the volume fraction of austenite. For this purpose, a Bruker™ D8 advance diffractometer was used with a CuKα radiation source operated at 40 kV and 40 mA. The scans were performed in the range of 35 to 110° (2θ) with a step size of 0.01° and a step time of 2 s according to the ASTM E975 standard. The Rietveld method was used for quantitative analysis of the XRD patterns. In this method, two models are generally used simultaneously for profile fitting of the diffraction patterns: (i) a structural model that takes into account the approximate position of the atoms, and (ii) a non-structural model that takes into account Bragg diffraction [18].

Analytical measurements were performed using Maud version 2.55 (2015) and HighScore Plus® version 3.0 programs to characterize the qualitative and quantitative features of the microstructures.

2.3 Mechanical testing

The Rockwell C hardness (HRC) test was performed using a ReicherterTM hardness tester in accordance with the ASTM E18 standard. A minimum of five measurements were taken and the average value was reported with an uncertainty of ±0.2 HRC within a 95% confidence level.

The CNT specimen is used to evaluate the fracture behavior of the steel as suggested by Alaneme [19]. The advantage of a CNT specimen compared to a conventional CT specimen is that a plane strain state can be achieved even at a smaller size [20]. Since the plastic zone in the plane strain deformation mode is smaller than that in plane stress, this leads to higher brittleness. Therefore, the fracture toughness in plane strain mode is investigated. For this purpose, the CNT specimens were machined to the dimensions shown in Figure 3, similar to what was previously described by Bayram et al. [21]. A tolerance of ±0.1 mm was maintained in the machining of the specimens according to the standard DIN 7168. The radius of curvature of each notch was controlled using a Leitz Wetzlar™ optical microscope. Tensile tests were performed using an InstronTM 8502 testing machine at a strain rate of 3 × 10−4 s−1.

Cajner et al. [6] showed that the fracture toughness of high-speed steels with precracked CNT specimens follows a linear behavior that can be predicted by equation (1). The same equation was later introduced by Alaneme in his studies on the fracture toughness of dual-phase medium carbon steel specimens without pre-cracking.(1)

where D is the outer diameter, d is the notch diameter, P is the maximum load in the load-displacement diagram, and KIc is the fracture toughness in MPa m1/2. Leskovsek and Liščić [12] have defined a new equation concerning the microstructural characteristics taking into account the factors of undissolved carbides and volume fraction of austenite phase in the fracture toughness equation:(2)where dp equals to:(3)where E is the modulus of elasticity in MPa, dp is the average distance between undissolved carbides in m, Dp is the average diameter of primary undissolved carbides in m, σ ys is the compressive yield stress in MPa, is the equivalent critical local fracture strain, f CARB and f AUS are the volume fractions of undissolved carbides and retained austenite, respectively. Therefore, K Ic_St is the calculated fracture toughness in MPa m1/2.

thumbnail Fig. 3

The geometry of the CNT sample. Dimensions are in mm with a tolerance of ±0.1 mm.

3 Results and discussion

3.1 Microstructure

The micrographs of marquenched and austempered microstructures are shown in Figures 4a and 4b, respectively, and show the bainitic and martensitic transformations before triple tempering (hereafter, tempering). The samples were deeply etched to reveal the prior austenite grain boundaries. In the marquenched sample, the matrix consists of plate martensite and austenite/martensite blocks, while in the austempered sample it consists of bainite within retained austenite and some plate martensite. In both samples, the undissolved carbides were uniformly distributed in the matrix. Prior austenite grain boundaries are clearly visible due to the deep etching. According to the CCT diagrams [2224], the martensite start temperature (Ms) of the steel is within 100–150 °C. Considering that the martensite finish temperature (MF) is about 200–220 °C lower than the Ms temperature [22], the occurrence of retained austenite in the marquenched samples is expected and justifiable. The selected austempering temperature of 235 °C is well above the Ms temperature to obtain a bainitic structure.

Grain boundaries (GB), undissolved carbides (UC), austenite (A), martensite (M) and bainite (B) laths are shown with arrows [23].

The microscopic images of both samples in Figures 4a and 4b show that the undissolved carbides are almost round and have an average diameter of about 1 µm. The uniformity in dispersion and roundness of the fine undissolved carbide particles can be attributed to the HIP process in the production of PM HSS ASP2030 [23]. The uniformity of the primary microstructure leads to more consistent test results. The acquired data of prior austenite grain size are summarized in Table 2 for samples quenched from austenitizing temperatures of 1150, 1160, 1170, and 1185 °C to room temperature (25 °C) without tempering. The distribution of ferrite laths makes it difficult to distinguish the prior austenite grain boundaries, as shown in Figure 4b. Moreover, low-temperature austempering (i.e. 235 °C) hinders the growth of austenite grains, as confirmed by Neumeyer and Kasak [24]. Consequently, only the grain size of prior austenites in quenched samples was investigated in this study.

The average diameter (Dp ), the volume fraction of undissolved carbides (fCARB ) and average distance between undissolved carbides (dp ) in marquenched and austempered samples after tempering are shown in Table 3.

From Table 3, the higher the austenitizing temperature, the greater the distance between undissolved carbides (dp ) and their size (Dp ) in the microstructure, resulting in a decrease in the volume fraction of carbides (fCARB ), as shown in Figure 5. A higher austenitizing temperature naturally leads to a greater dissolution of the alloying elements, so that the volume fraction of carbides decreases and carbides with coarser size and greater spacing between them are formed, in accordance with the phenomenon of Ostwald ripening [25].

thumbnail Fig. 4

FE-SEM images of samples before tempering a) M8–1150 b) A60–1170.

Table 2

Maximum inscribed and circumscribed circle diameters in and out of the prior austenite grains.

Table 3

The volume fraction and the mean diameter of carbides after tempering for marquenched and austempered samples austenitized at different temperatures.

thumbnail Fig. 5

Variation of the volume fraction of undissolved carbides with austenitizing temperature for marquenched and austempered samples.

3.2 X-ray diffraction

Figure 6 shows an XRD pattern of the heat-treated M8-1160 sample after tempering where the fitted and residual curves also shown. The nearly horizontal residual curve indicates good refinement by the Rietveld method [18,26]. The pattern of the tempered sample of M8-1160 shows the peaks of ferrite, austenite and various types of carbides such as M6C like (Co3W3)C or Fe3(W, Mo)3C, and V8C7 [27].

The volume fractions of austenite and ferrite in the marquenched and austempered samples after triple tempering are calculated and summarized in Table 4. The austenite content in the tempered samples is about 5–6% by volume. The low content of retained austenite indicates that the tempering process is completed [28]. The degree of fit was evaluated by minimizing the weighted profile R-factor (Rwp) and the Rietveld standard deviations of the structural parameters [29].

According to Tated and Patil [14], the wear resistance of hardened and tempered HSS increases with the decrease of the austenite content and the formation of carbides.

Inhomogeneous plastic deformation can easily produce induced stresses in the crystal, which increase the microstrains within the crystal. The microstrain usually broadens the XRD peaks [30]. The microstrain ε is given below [31]:(4)

where θ is the Bragg angle of the diffraction peak and βhkl is the full-width at half-maximum (FWHM) of the Bragg’s reflection of (hkl) planes corresponding to the ferrite phase and determined by equation (5) [31]:(5)

The net peak area was determined by the Rietveld method by fitting a pseudo Pseudo-Voigt function to the XRD data using HighScore plus® software.

The crystallite sizes can also be determined using the Scherrer equation (6) [32].(6)where k is the shape factor and is equal to 0.9 and λ is the wavelength of the CuKα radiation.

The peaks of the (110) and (211) planes were used for the calculations because the (200) planes showed interference with the carbide peaks, see Figure 7. The non-uniform stress fields due to the supersaturation of the martensite by carbon atoms led to the accumulation of dislocations, which caused the broadening of the diffraction peaks [33]. Also, the ferrite peaks were shifted to lower diffraction angles, resulting in smaller plane spacings.

The calculated values for βhkl, interplanar spacing dhkl, microstrain, crystallite size and goodness of fitting ratio (χ 2) often used in the literature [18,26] are tabulated in Table 5.

The results in Table 5 show that with increasing austenitizing temperature, both FWHM and microstrain increased in the marquenched and austempered samples, which seems to be due to the higher amount of dissolved alloying elements, including carbon, in the body-centred tetragonal (BCT) lattice. The decrease in FWHM at 1185 °C is due to the larger increase in crystal size, as shown in Table 5.

thumbnail Fig. 6

The X-ray diffraction pattern of M8–1160 after triple tempering indicated by experimental points, the calculated fitted curve and the residual curve.

Table 4

The ferrite and austenite volume fractions and weighted profile R factor in austempered and marquenched samples calculated by the Rietveld method.

thumbnail Fig. 7

Shape profile fitted on diffraction pattern of M8–1170 after triple tempering.

Table 5

Diffraction angle, full-width at half-maximum (βhkl) of the Bragg reflection and the goodness of fitting ratio (χ 2) for the raw material and heat-treated samples.

3.3 Hardness

The hardness values of the heat treated and tempered samples for different austenitizing temperatures are summarized in Table 6.

Although the hardness variation with austenitizing temperature is small, the slight increase in hardness due to the increase in austenitizing temperature from 1150 to 1170 °C may be related to the greater dissolution of alloying elements in austenite [25], which leads to an increase in the volume fraction of fine secondary carbides with dimensions of 10–30 nm × 4–5 nm (length × thickness) formed during tempering treatment [2,34]. On the other hand, the measured microstrain shows the lowest values at the highest austenitizing temperature (Tab. 5). The slight decrease in hardness at the austenitizing temperature of 1185 °C can be explained by the combined effect of microstrain and alloy content of austenite. This is consistent with the argument already made by Galindo-Nava and Rivera-Díaz-del-Castillo [35] that the increase in hardness with increasing dislocation density of martensitic laths is directly related to the squared value of the lattice strain.

Table 6

Hardness values in heat-treated tempered samples for various austenitizing temperatures (Standard Deviation = ±0.2 RC).

3.4 Fracture toughness

The fracture toughness of the heat-treated CNT samples was calculated using equation (1), and the results are shown in Table 7.

The theoretical prediction of fracture toughness with microstructural and mechanical parameters was developed by Leskovsek and Liščić for M2 tool steels. The validity of the model, see equation (2), was investigated by considering the role of Rockwell hardness in the calculation of the equivalent critical local elongation at fracture strain (). It is assumed that is inversely proportional to the hardness of the matrix [12]:(7)

in equation (7), k1 is a constant and HRC is the Rockwell C hardness.

Then equation (2) can be rewritten as:(8)where the fracture toughness is derived from the mechanical response and microstructural parameters of the tool steel. The variation in hardness between all the heat treated samples ranged from 1 to 2 HRC. Therefore, can be taken as a constant value and the fracture toughness, KIc, is expressed as follows.(9)

Moreover, XRD analysis revealed that the volume fraction of retained austenite is low (i.e., fAUS  = 0.05–0.10) for all heat-treated samples after tempering. Therefore, it can be assumed that the small amount of retained austenite does not have a significant effect on the calculations:(10)

It is known that the yield stress of the matrix YS, is the sum of the different strengthening mechanisms [36]:(11)

Δσ 0Δσ 0 is the intrinsic strength of the matrix, which is constant, and ΔσSS is the strength increase due to the solid solution. Due to the low solubility of carbon in the high temperature tempered martensite matrix, the contributions of carbon and carbide forming elements are considered negligible. Δσ GB is the grain boundary strengthening and Δσ DIS is the dislocation strengthening [36]. Taking into account the equivalent expressions, Considering the equivalent expressions, equation (10) can be expressed as follows:(12)As mentioned above, it can be assumed that the intrinsic and the solid-solution strengthening mechanisms are constant over the heat treatment, then:(13)where d is the grain diameter and for the lath martensite the grain boundary effect has been determined by either the lath width or the package packet size, M is the Taylor factor, K HP and K α are constants, G is the shear modulus, b is the Burgers vector and ρ is the average dislocation density.

One can split the fracture toughness relation into two components, namely the unresolved carbide and microstructural part as , which contributes to the final fracture toughness calculation.

Then:(14)then:(15)

The dislocation density (ρ) is directly related to the lattice strain according to the equation (16) [32].(16)where <ϵ 2 > 1/2 is the square root of strain and D is the average crystallite size of these XRD peaks.

The calculated K [carb] based on microstructural features, microstrain (ε) and dislocation density (ρ) are listed in Table 8.

The fracture toughness of the marquenched and austenitized samples, KIc, increases with increasing austenitization temperature from 1150 to 1170 °C (Fig. 8). As the austenitization temperature increases, the spacing between the undissolved carbides (dP) increases while their volume fraction decreases (Tab. 3). The increase in dP and decrease in volume of the undissolved carbides leads to an increase in K [carb], which consequently increases KIc_St, see equation (14). This agrees well with the results of Blaha et al. [36] on tool steels, who showed that the large distance between particles produces a longer crack propagation path in the matrix, leading to fracture. In parallel, the fracture toughness is enhanced by the increase in the dislocation density as shown in equation (15), rearranged from equation (2). The calculated results listed in Table 8 reveals that the change in the fracture toughness is dictated by the simultaneous effect of microstructural components shown with the parameter with the austenitization temperature up to 1170 °C. The decrease in the fracture toughness at the austenitizing temperature of 1185°C for marquenched samples can be explained by parameter although K[carb] at this temperature has a higher value in comparison with the austenitizing temperature of 1150 and 1160 °C. Due to further decrease in dislocation density at 1185 °C compared to lower austenitizing temperatures, K  [carb] and consequently, the fracture toughness decreases. A similar effect is observed for austempered samples about the influence of austenitization temperature on fracture toughness.

The KIc of the austempered samples at all austenitizing temperatures is higher than those of marquenched ones, see Figure 8. Crane and Bigg [5] also reported that changing the austenitization and tempering temperature of AISI M2, T1 high-speed steel results in an increase in fracture toughness of about 5 MPa m1/2, while the improvement due to austenitization is about 10 MPa m1/2.

The reason can be explained from Table 8 which the austempered samples have a higher value of the parameter than the Marquenched samples.

Figure 9 shows the simultaneous influence of K [carb] and dislocation density in the form of an interpolated contour map and 3D scatter plat, processed using Matlab® software based on the data reported in Table 8. For higher values of fracture toughness, the simultaneous increase of both K [carb] and dislocation density are necessary.

The plastic zone radius, ry , under the plane strain condition can be calculated by the following formula [37]:(17)

Using the equation (17), the calculated plastic zone radii for several heat-treated samples are listed in Table 9. Since the values of ry are almost equal to or less than the prior austenite grains radius, Table 2, then the grains size effect on fracture toughness was not significant for the studied tool steel.

Naylor [38] has reported that a decrease in either lath width or packet diameter of lath type martensitic or bainitic steels simultaneously increase both yield stress and fracture toughness of the steel by hindering the dislocation glide through lath boundaries.

EBSD analysis was used for the evaluation of phases and to measure their thickness distribution more precisely. The images were taken for A60–1150 and M8–1170 samples indicate a lath type microstructure for both samples (Fig. 10).

The thickness distribution of laths shown in Figure 11 indicates that the ferrite laths thickness for both marquenched and austempered samples is smaller than the calculated ry , then it can be considered as a factor in the variation of fracture toughness. According to Table 8 the value of is lower for A60–1150 than for the M8–1170 sample. Then it is predicted that the fracture toughness of A60–1150 ought to be higher than that of the M8–1170 sample. But the results show that the fracture toughness of the A60–1150 (20.7 MPa m1/2) sample is higher than that of the M8–1170 (16.7 MPa m1/2) sample. This contradiction can be explained by the observed results for the corresponding microstructure’s refinement. The bainite laths limit the growth of martensite plates and contribute to the refinement of the final microstructure [17,39]. A similar result was observed by Zhao et al. [40], who studied a medium carbon, high silicon steel with low temperature austemperig,and reported that the strength and fracture toughness increased due to the finer bainitic ferrite laths.

The impact of lath size (d) is also considered in equation (12), where the fracture toughness improvement depends on the yield stress increase as a result of lath size (d) reduction. Furthermore, according to XRD results, with decreasing the crystallite size (D), the dislocations density increases according to equation (16), resulting in the fracture toughness improvement. Consequently, apart from the dislocation density, undissolved carbides size and volume, the laths average thickness has an important role in the fracture toughness of tool steels as a microstructural feature. Since the austenite volume is relatively low according to XRD results and the EBSD analysis did not reveal it, then its influence on the fracture toughness was ignored in the present study.

Table 7

Fracture toughness of heat-treated samples after tempering.

Table 8

Calculated K [carb] and microstrain values of heat-treated samples.

thumbnail Fig. 8

KIc values for heat-treated samples.

thumbnail Fig. 9

The simultaneous influence of dislocation density and K[carb] on fracture toughness in the form of (a) interpolated contour map and (b) 3D scatter plot.

Table 9

Comparison of the plastic zone radius with microstructure specifications.

thumbnail Fig. 10

EBSD image quality maps for (a) M8–1170 (undissolved carbides, UC, shown with arrows) (b) A60–1150 and inverse pole figures & boundary Maps for (c) M8–1170 and d) A60–1150 samples.

thumbnail Fig. 11

The thickness distribution of ferrite laths in the matrix of (a) A60–1150 and (b) M8–1170 samples.

4 Conclusions

By studying the effects of microstructural features on the fracture toughness of austempered and marquenched PM HSS ASP2030 austenitized at different temperatures, the following conclusions can be drawn:

  • In both heat treatment conditions, the amount of retained austenite was small and almost the same, indicating that the spacing between carbides, the volume fraction of undissolved carbides as microstructural features, and the dislocation density together with the size of the laths as nanostructural features have a significant influence on the fracture toughness.

  • By decreasing the volume fraction of undissolved carbides, the distance between carbides increases and the fracture toughness increases. Also, by increasing the dislocation density resulting from microdeformation of the lattice, the fracture toughness and hardness increase. The simultaneous effect of volume fraction of undissolved carbides and lattice microstrain can explain the change in fracture toughness and hardness at different austenitizing temperatures.

  • The fracture toughness of austempered microstructure was higher than that of the marquenched microstructure because it has a larger spacing between carbides, a lower content of undissolved carbides, and thinner bainitic-martensitic laths than the marquenched microstructure.

Acknowledgements

The authors would like to thank Sahand University of Technology for providing the research facilities and financial support.

References

  1. A.M. Bayer, B.A. Becherer, T. Vasco, High Speed Tool Steels, ASM Handbook, 1989, vol. 16, pp. 51–59, www.asminternational.org [Google Scholar]
  2. W. Rongi, H.O. Andren, H. Wisell et al., The role of alloy composition in the precipitation behavior of high speed steel, Acta Metall. Mater. 40, 1727–1738 (1992) [CrossRef] [Google Scholar]
  3. S. Lee, C.G. Lee, K.-S. Sohn et al., Correlation of microstructure and fracture toughness in three high-speed steel rolls, Metall. Mater. Trans. A 28, 123–134 (1997) [CrossRef] [Google Scholar]
  4. N. Sarafianos, The effect of the austenitizing heat-treatment variables on the fracture toughness of high-speed steel, Metall. Mater. Trans. A 28, 2089–2099 (1997) [CrossRef] [Google Scholar]
  5. L.W. Crane, A.P. Bigg, Fracture toughness of high speed steels, Mater. Sci. Technol. 6, 993–998 (1990) [CrossRef] [Google Scholar]
  6. F. Cajner, V. Leskovšek, D. Landek et al., Effect of deep-cryogenic treatment on high speed steel properties, Mater. Manufactur. Process. 24, 743–746 (2009) [CrossRef] [Google Scholar]
  7. M.Y. Tu, C.A. Hsu, W.H. Wang et al., Comparison of microstructure and mechanical behavior of lower bainite and tempered martensite in JIS SK5 steel, Mater. Chem. Phys. 107, 418–425 (2008) [CrossRef] [Google Scholar]
  8. T. Teker, S.O. Yilmaz, Investigation on mechanical properties of solution strengthened and austempered ferritic ductile irons, Int. J. Mater. Res. 111, 976–982 (2020) [CrossRef] [Google Scholar]
  9. J.A. Kaleicheva, Structure and properties of high-speed steels after austempering, Int. J. Microstruct. Mater. Proper. 2, 16–23 (2007) [Google Scholar]
  10. Y. Xuejie, Z. Tianjian, Secondary hardening of bainite in high speed steel, Acta. Metall. Sin. −Engl. 3, 99–103 (1990) [Google Scholar]
  11. A. Güral, O. Altuntaş, Improving the impact toughness properties of high carbon powder metallurgy steels with novel spherical cementite in the bainitic matrix (SCBM) microstructures, Mater. Chem. Phys. 259, 124203 (2021) [CrossRef] [Google Scholar]
  12. B.U.V. Leskovsek, B. Liščić, Relations between fracture toughness, hardness and mircrostructure of vacuum heat-treatment high-speed steel, J. Mater. Process. Technol. 127, 298–308 (2002) [CrossRef] [Google Scholar]
  13. ERASTEEL, ASP2030. [Online]. Available: https://www.erasteel.com [Google Scholar]
  14. P.I. Patil, R.G. Tated, Comparison of effects of cryogenic treatment on different types of steels: a review, in International Conference in Computational Intelligence (ICCIA), Taylor & Francis Online, 2012, vol. 9, pp. 10–29 [Google Scholar]
  15. R. Vilar, R. Colaço, A. Almeida, Laser surface treatment of tool steels, Laser Process. 307, 453–478 (1996) [Google Scholar]
  16. B. Shakerifard, J.G. Lopez, L.A.I. Kestens, A new electron backscatter diffraction-based method to study the role of crystallographic orientation in ductile damage initiation, Metals 10, 113 (2020) [CrossRef] [Google Scholar]
  17. F. Wang, D. Qian, H. Mao et al., Evolution of microstructure and mechanical properties during tempering of M50 steel with Bainite/Martensite duplex structure, J. Mater. Res. Technol. 9, 6712–6722 (2020) [CrossRef] [Google Scholar]
  18. G. Will, Powder diffraction: the rietveld method and the two stage method to determine and refine crystal structures from powder diffraction data, Berlin Heidelberg, Springer-Verlag, 2006 [Google Scholar]
  19. K. Alaneme, Fracture toughness (K-IC) evaluation for dual phase medium carbon low alloy steels using circumferential notched tensile (CNT) specimens, Mater. Res. 14, 155–160 (2011) [CrossRef] [Google Scholar]
  20. A.F. Liu, Mechanics and mechanisms of fracture: an introduction, ASM International, 2005 [CrossRef] [Google Scholar]
  21. A. Bayram, A. Uguz, Rapid determination of the fracture toughness of metallic materials using circumferentially notched bars, J. Mater. Eng. Perform. 11, 571–576 (2002) [CrossRef] [Google Scholar]
  22. N.A.D. Candane, K. Palaniradja, Effect of cryogenic treatment on microstructure and wear characteristics of AISI M35 HSS, Int. J. Mater. Sci. Appl. 2, 56–65 (2013) [Google Scholar]
  23. H. Takigawa, H. Manto, N. Kawai et al., Properties of high-speed steels produced by powder metallurgy, Powder Metall. 24, 196–202 (1981) [CrossRef] [Google Scholar]
  24. T.A. Neumeyer, A. Kasak, Grain size of high-speed tool steels, Metall. Trans. B, 2281–2287 (1972) [CrossRef] [Google Scholar]
  25. Y. Luo, H. Guo, X. Sun et al., Effects of austenitizing conditions on the microstructure of AISI M42 high-speed steel, Metals 7, 1–13 (2017) [Google Scholar]
  26. L. McCusker, R. Von Dreele, D. Cox et al., Rietveld refinement guidelines, J. Appl. Crystallogr. 32, 36–50 (1999) [CrossRef] [Google Scholar]
  27. M. Wieβner, M. Leisch, H. Emminger et al., Phase transformation study of a high speed steel powder by high temperature X-ray diffraction, Mater. Charact. 59, 937–943 (2008) [CrossRef] [Google Scholar]
  28. B. Liu, T. Qin, W. Xu et al., Effect of tempering conditions on secondary hardening of carbides and retained austenite in spray-formed M42 high-speed steel, Materials 12, 3714 (2019) [CrossRef] [Google Scholar]
  29. S. Takebayashi, T. Kunieda, N. Yoshinaga et al., Comparison of the dislocation density in martensitic steels evaluated by some X-ray diffraction methods, ISIJ Int. 50, 875–882 (2010) [CrossRef] [Google Scholar]
  30. M. Müller, M. Zehetbauer, A. Borbely et al., Dislocation density and long range internal stresses in heavily cold worked Cu measured by X-ray line broadening, Int. J. Mater. Res. 86, 827–831 (1995) [CrossRef] [Google Scholar]
  31. Y.-J. Shen, Y.-L. Zhang, F. Gao et al., Influence of temperature on the microstructure deterioration of sandstone, Energies 11, 1–17 (2018) [Google Scholar]
  32. M.S.K. Venkateswarlu, T.A. Nellaippan, N. Rameshbabu, Estimation of crystallite size, lattice strain and dislocation desity of nanocrystalline carbonate substituted hydroxyapatite by X-ray peak variance analysis, Proc. Mater. Sci. 5, 212–221 (2014) [CrossRef] [Google Scholar]
  33. N. Du, H. Liu, P. Fu et al., Microstructural stability and softening resistance of a novel hot-work die steel, Crystals 10, 238 (2020) [CrossRef] [Google Scholar]
  34. W. Bochnowski, H. Leitner, R. Ebner et al., Primary and secondary carbides in high-speed steels after conventional heat treatment and laser modification, Mater. Chem. Phys. 81, 503–506 (2003) [CrossRef] [Google Scholar]
  35. E. Galindo-Nava, P. Rivera-Díaz-del-Castillo, Understanding the factors controlling the hardness in martensitic steels, Scr. Mater. 110, 96–100 (2016) [CrossRef] [Google Scholar]
  36. H. Liu, P. Fu, H. Liu et al., Improvement of strength-toughness-hardness balance in large cross-section 718H pre-hardened mold steel, Materials 11, 1–19 (2018) [Google Scholar]
  37. A.F. Liu, Mechanics and mechanisms of fracture: An introduction (ASM International,the Unit ed States of America, Ohio). ASM International, 2005 [CrossRef] [Google Scholar]
  38. J.P. Naylor, The influence of the lath morphology on the yield stress and transition temperature of martensitic-bainitic steels, Metall. Trans. A 10, 861–873 (1979) [CrossRef] [Google Scholar]
  39. Y. Ohmori, H. Ohtani, T. Kunitake, Tempering of the bainite and the bainite/martensite duplex structure in a low-carbon low-alloy steel, Met. Sci. J. 8, 357–366 (1974) [CrossRef] [Google Scholar]
  40. T. Zhao, X. Jia, C. Chen et al., Introducing ultrafine ferrite in low-temperature bainitic steel through a novel process for simultaneously improving strength and toughness, J. Mater. Res. Technol. 15, 5106–5113 (2021) [CrossRef] [Google Scholar]

Cite this article as: Ahmad Firouzi, Sasan Yazdani, Reza Tavangar, Behnam Shakerifard, Faseeulla Khan Mohammad, Austempering of PM HSS ASP2030 for improved fracture toughness, Metall. Res. Technol. 119, 211 (2022)

All Tables

Table 1

Chemical composition of PM HSS ASP2030 (wt.%).

Table 2

Maximum inscribed and circumscribed circle diameters in and out of the prior austenite grains.

Table 3

The volume fraction and the mean diameter of carbides after tempering for marquenched and austempered samples austenitized at different temperatures.

Table 4

The ferrite and austenite volume fractions and weighted profile R factor in austempered and marquenched samples calculated by the Rietveld method.

Table 5

Diffraction angle, full-width at half-maximum (βhkl) of the Bragg reflection and the goodness of fitting ratio (χ 2) for the raw material and heat-treated samples.

Table 6

Hardness values in heat-treated tempered samples for various austenitizing temperatures (Standard Deviation = ±0.2 RC).

Table 7

Fracture toughness of heat-treated samples after tempering.

Table 8

Calculated K [carb] and microstrain values of heat-treated samples.

Table 9

Comparison of the plastic zone radius with microstructure specifications.

All Figures

thumbnail Fig. 1

Dilatometry curve to measure the AC1 and AC3 temperatures.

In the text
thumbnail Fig. 2

Schematic illustration of austempering and marquenching heat treatment cycles.

In the text
thumbnail Fig. 3

The geometry of the CNT sample. Dimensions are in mm with a tolerance of ±0.1 mm.

In the text
thumbnail Fig. 4

FE-SEM images of samples before tempering a) M8–1150 b) A60–1170.

In the text
thumbnail Fig. 5

Variation of the volume fraction of undissolved carbides with austenitizing temperature for marquenched and austempered samples.

In the text
thumbnail Fig. 6

The X-ray diffraction pattern of M8–1160 after triple tempering indicated by experimental points, the calculated fitted curve and the residual curve.

In the text
thumbnail Fig. 7

Shape profile fitted on diffraction pattern of M8–1170 after triple tempering.

In the text
thumbnail Fig. 8

KIc values for heat-treated samples.

In the text
thumbnail Fig. 9

The simultaneous influence of dislocation density and K[carb] on fracture toughness in the form of (a) interpolated contour map and (b) 3D scatter plot.

In the text
thumbnail Fig. 10

EBSD image quality maps for (a) M8–1170 (undissolved carbides, UC, shown with arrows) (b) A60–1150 and inverse pole figures & boundary Maps for (c) M8–1170 and d) A60–1150 samples.

In the text
thumbnail Fig. 11

The thickness distribution of ferrite laths in the matrix of (a) A60–1150 and (b) M8–1170 samples.

In the text

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