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Improved the microstructures and properties of M3:2 high-speed steel by spray forming and niobium alloying
Materials Characterization 117 (2016) 1–8
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Improved the microstructures and properties of M3:2 high-speed steel by spray forming and niobium alloying L. Lu a, L.G. Hou a,⁎, J.X. Zhang a, H.B. Wang a, H. Cui a, J.F. Huang a, Y.A. Zhang b, J.S. Zhang a a b
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Xueyuan Road 30, Haidian District, Beijing 100083, China State Key Laboratory of Non-Ferrous Metals and Process, General Research Institute for Non-Ferrous Metals, Beijing 100088, China
a r t i c l e
i n f o
Article history: Received 21 May 2015 Received in revised form 16 April 2016 Accepted 19 April 2016 Available online 20 April 2016 Keywords: High-speed steel Spray forming Mechanical properties Nb alloying
a b s t r a c t The microstructures and properties of spray formed (SF) high-speed steels (HSSs) with or without niobium (Nb) addition were studied. Particular emphasis was placed on the effect of Nb on the solidification microstructures, decomposition of M2C carbides, thermal stability and mechanical properties. The results show that spray forming can refine the cell size of eutectic carbides due to the rapid cooling effect during atomization. With Nb addition, further refinement of the eutectic carbides and primary austenite grains are obtained. Moreover, the Nb addition can accelerate the decomposition of M2C carbides and increase the thermal stability of high-speed steel, and also can improve the hardness and bending strength with slightly decrease the impact toughness. The high-speed steel made by spray forming and Nb alloying can give a better tool performance compared with powder metallurgy M3:2 and commercial AISI M2 high-speed steels. © 2016 Elsevier Inc. All rights reserved.
1. Introduction High-speed steels (HSSs), as a special class of highly alloyed tool steels, are extensively used as cutting tools, cold molds and rolls for the exceptional combination of excellent hot hardness, high wear resistance, well toughness and good workability. These properties are mainly depended on the size, type, distribution, shape and amount of carbides and the state of the matrix [1–3]. HSSs made by conventional ingot casting contain coarse network of eutectic carbides formed at grain boundaries. Even after certain forging degree, the carbide particles are arranged in bands parallel to forging direction, which deteriorating the isotropic properties and toughness of HSSs [4,5]. On the other hand, the slow cooling rate of conventional ingot casting limited the addition of strong carbide forming elements such as Ti, Nb, V due to the formation of coarse carbides. Fortunately, the rapid solidification technology provides a new opportunity to modify the solidification microstructures and improve the properties of HSSs, e.g., powder metallurgy (PM) was the first and the most mature industrial application of the rapid solidification technology for highspeed steel products. Finer and more uniform primary carbides, smaller grain sizes and the absence of carbides stringers and macro-segregation are some characteristics attained by PM HSSs compare with the conventional ingot casting HSSs [6–8]. Also it is possible to produce any alloy design by PM process. However, the PM HSSs is still applied in a narrow field because they have normally high cost due to the complex and rigorous processing steps. ⁎ Corresponding author at: University of Science and Technology Beijing, Beijing, China. E-mail address: [email protected] (L.G. Hou).
http://dx.doi.org/10.1016/j.matchar.2016.04.010 1044-5803/© 2016 Elsevier Inc. All rights reserved.
Spray forming (SF), a promising preparation technology for materials, can combine the rapid solidification (gas atomization) and nearnet-shape forming (deposition) and has been proved to be an effective method to improve the solidification microstructures of tool steels [9–12], due to the high cooling rate (103–105 K/s) during gas atomization and the microstructural deformation or fracture in deposition stage [13]. In contrast to conventional ingot casting, the most significant microstructure characteristics of SF HSSs are the smaller equiaxed grains, finer and much more uniform-distributed carbides, and less degree of segregation [6,8], resulting in the improved performance of SF HSSs. However, the increase of cooling rate in atomization stage is limited by the density of the final billet [14–16]. Alloying design has been widely used to improve the microstructures and properties of HSSs, and two main alloying strategies are proposed: adding non-carbide forming elements (e.g., Co, Al, Si and rare earth) [17–19], or modifying the microstructures and properties with strong carbide forming elements (e.g., Ti, Nb, Zr). Recently, a relatively great number of literatures reflect the enhanced interest of researchers in various countries on the use of Nb and Ti for alloying HSSs [20–27], but most of the mentioned works were based on conventional ingot casting. In this paper, the idea of combining spray forming and Nb alloying to improve the performance of M3:2 HSS is implemented and the purpose of the present work is to study the microstructures and mechanical properties of spray formed Nb-containing HSSs as well as the role of Nb in HSSs. 2. Experimental procedure Billet with dimension of Φ180 mm × 100 mm was prepared by spray forming and the processing parameters are shown in Table 1. A
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L. Lu et al. / Materials Characterization 117 (2016) 1–8 Table 1 Process parameters of spray forming. Items
Parameters
Super heat/°C Diameter of the delivery tube/mm Atomization gas Atomization pressure/MPa Deposition distance/mm
160–180 4.0 N2 0.45–0.5 425–450
Table 2 Chemical compositions of the investigated steels (wt.%).
Steel A Steel B
C
W
Mo
Cr
V
Nb
1.30 1.31
6.20 6.10
5.10 4.90
4.60 4.48
2.80 2.75
– 0.50
commercial AISI M2 HSS used as the feedstock was melted with appropriate additions of alloying elements. The chemical compositions of the experimental steels are listed in Table 2. A cylindrical specimen with a size of Φ50 mm × 40 mm was cut from the center of the spray-formed billet and subjected to isothermal annealing at 1180 °C for 1.5 h, then forged to Φ16 mm bars at 1150 °C followed with slowly cooling to room temperature. Before the hardening treatment, all the specimens were heated to 900 °C with a heating rate of 100 °C/h and soaked for 1.5 h, then cooled to 25 °C with a cooling rate of 20 °C/h. Subsequently, the specimens were austenitized at 1160–1220 °C for 25 min and oil quenched to room temperature, followed by triple tempering at 540–600 °C for 1 h. In order to prevent decarburization during heat treatment, each sample was sealed in vacuum quartz tube. The microstructures were characterized using field emission scanning electron microscope (SEM), and Energy dispersive X-ray spectroscopy (EDS) was employed for local chemical composition determination and for the distinction of MC and M2C carbides. For metallographic examinations, the samples were prepared by grinding and polishing, and were etched with 8% nital solution. The carbide volume fraction as well as grain size was measured by using image analysis software (Image-pro plus 6). The reported values for each sample were
taken from at least 50 measurements in different fields with a magnification of 1000 times. The extraction of carbides was performed using electrolysis, operating at 40 V, 0 °C with a electrolyte contained 7 g citric acid, 20 mL hydrochloric acid and 250 mL methanol. X-ray scattering was made with a Rigaku X-ray Diffractometer (D/MAX-RB) with CuKα radiation (λ = 1.5406 Å) at 12 kW. Differential scanning calorimetry (DSC) experiments were carried out in a Netzsch DSC calorimeter with cooling rate of 10 K/min from 1455 to 200 °C. The DSC samples (Φ5 mm × 1 mm) were machined from the asdeposited billet. Three-point bending tests were carried out on specimens with dimension of 5 × 5 × 35 mm at a bending rate of 0.1 mm/min. The sample size for impact toughness tests was 10 × 10 × 55 mm without notch according to ISO 5754. The bulk hardness measurements were carried out using Rockwell “C” scale with a load of 150 kgf (589 N) applied for 3 s and 7–10 readings were considered for estimating the hardness values. The reamers machined from the forged bars were sealed in vacuum quartz tube and annealed at 900 °C for 1.5 h, and then austenitized at 1200 °C for 25 min and oil-quenched to room temperature, followed by triple tempering at 560 °C for 1 h. The reamers used to ream the medium carbon steel with average hardness of 32HRC, and all the tests were under the condition of oil lubrication. 3. Results and discussion 3.1. Effects of spray forming and Nb alloying on the microstructures of HSSs The microstructures of the conventional ingot casting HSS and SF HSS in Fig. 1 show that the solidification microstructures are consisted of equiaxed grains, discontinuous network of plate-shaped M2C eutectic carbides (white) and MC carbides (gray) both at grain boundaries and within the matrix. It can be obtained that the grain size of SF HSS is similar to that of conventional ingot casting HSS (~60 μm). However, the sizes of M2C and MC carbides in SF HSS are remarkably refined (Fig. 1a and b), e.g., the cell sizes of M2C eutectic carbides in conventional ingot casting HSS and SF HSS are ~35 μm, ~17 μm, respectively. R.A. Mesquita and C.A. Barbosa [4] reported that high cooling rate of atomization is the main factor for primary carbides refining. Thus, the cooling rate could be the key factor to refine the solidification microstructure of
Fig. 1. Microstructures of steel A (a, b) and B (c) made by ingot cast (a) and spray forming (b, c).
L. Lu et al. / Materials Characterization 117 (2016) 1–8
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particles in steel B can be act as the inoculants for the nucleation during eutectic reaction and refine the eutectic carbides. These two factors are the main reason for the refinement of M2C carbides. From the composition of MC and M2C carbides in the spray-formed HSSs (Table 3), it is can be found that Nb alloying can decrease the contents of W, V while increase the content of Fe in M2C carbides, which will be beneficial to the decomposition of M2C carbides (in Section 3.2). Also the composition of MC carbides is changed remarkably with Nb addition, about 25% of V is substituted by 18% of Nb, and the content of W and Mo are increased. 3.2. Effect of Nb on the decomposition of M2C carbides
SF HSSs. But much higher cooling rates may decrease the liquid fraction of atomized particles and eventually result in the density reduction of the as-sprayed billet. In order to further refine the as-deposited microstructures, Nb is added and Fig. 1c shows excellent microstructures which can be characterized as fine grains, uniform distribution of fine M2C eutectic carbides and MC carbides. The mean cell sizes of M2C eutectic carbides in Fig. 1c is ~9 μm and the grain size is 45 μm. The volume fraction of M2C eutectic carbides decreases from 5.33% to 4.05% and the volume fraction of MC carbides increases from 2.68% to 3.12% with Nb addition (Fig. 1b and c). Generally, it can be concluded that the spray-formed HSSs can have much finer microstructures and less degree of carbides segregation than that of conventional ingot casting HSSs. Adding Nb also can considerably refine the prior-austenite grain sizes, reduce the cell size and volume fractions of M2C eutectic carbides while increase the amounts of MC carbides. Thus, the spray forming process combining with Nb alloying can effectively modify the solidified microstructures of HSSs. Basing on the results of ref [28,29], the solidification sequences of M3:2 high-speed steel with high carbon content can be described as follows:
M2C eutectic carbide which is a meta-stable phase in as-deposited microstructures will deteriorate the hot workability and toughness [5]. According to practical experience we expect that it has been decomposed prior to hot working. There are some reports [30–33] indicate that M2C carbides will decompose into M6C and MC carbides when annealed at high temperatures (N1000 °C). It is conceivable that using the decomposition treatment to obtain a more uniform carbides size and distribution, indeed to improve the performance of HSSs. To explore this possibility, it is necessary to study the effect of Nb on the decomposition of M2C carbides. Fig. 3 shows the microstructures of spray formed steels annealed at different temperatures for 1 h. At 1000 °C, no decomposition of M2C carbides can be found in steel A (Fig. 3a) and steel B (Fig. 3b). Increasing the annealing temperature to 1050 °C, most of the M2C carbides in steel A are still not decomposed except a few of small sized M2C carbides distributed on the edge of eutectic cells (arrowed in Fig. 3c). While all of the M2C carbides in steel B are decomposed into M6C and MC carbides, and the MC carbides are distributed as small dots inside the M6C carbides (Fig. 3d). But the new carbides still keep the original plate-shape of M2C carbides. At 1100 °C, partial decomposition of large sized M2C carbides in steel A (Fig. 3e) are obtained. But in steel B, the separation occurs at the interface between M6C and MC carbides, and the necking areas are formed in the thinner regions of the plates (Fig. 3f). The XRD results of the extracted carbides (Fig. 4) also show the peaks of M2C carbides in steel A, no peaks of M2C carbides were detected in steel B, meaning the complete decomposition of M2C carbides at 1100 °C in steel B. The decomposition reaction of M2C carbides obeys the formula as follows [31]:
1. Precipitation of austenite from liquid.
M2 C þ γ→M 6 C þ MC:
Fig. 2. DSC curves for investigated steels cooled from 1450 °C.
2. Formation of carbides at the interface of γ and liquid. 3. Formation of eutectic carbides through eutectic reaction.
ð1Þ
H. Fredriksson [33] observed that the nucleation of M6C carbides occurred firstly at M2C/austenite interfaces, and the transformation would be proceeded with the nucleation of MC carbides at either M6C/austenite or M2C/M6C interfaces, and the decomposition rate was proposed to depend upon the diffusion of alloying elements. Mainly, M6C carbides need Fe atoms from the austenite and Mo/W atoms from M2C. At a constant decomposition treating temperature and time, the finer M2C plates will be decomposed quickly than the coarse ones. As mentioned in Section 3.1, Nb alloying can refine the M2C carbides, so the decomposition of M2C carbides in steel B could be faster than that in steel A. Moreover, the stability of M2C carbides in Nb-alloying steels will be
The DSC curves in Fig. 2 show that the presence of 0.5% Nb in M3:2 HSS will not change the solidification sequences, but it apparently increases the temperature for MC formation and decreases the liquidus temperature. That is to say the primary austenite precipitates at a lower temperature in steel B, thus the Nb containing steel has a finer austenite grains than Nb free steel (see Fig. 1b and c). Because of high affinity with carbon, Nb reacts with free carbon of the melt, decreasing the carbon content of the melt. In such a case, the amount of residual interdendritic liquid for eutectic reaction will decrease, and more MC
Table 3 Composition of primary carbides and the matrix (at.%). Steel
Carbides
W
Mo
Cr
V
Nb
Fe
A
M2C MC Matrix M2C MC Matrix
39.02 ± 0.54 11.94 ± 0.61 1.48 ± 0.69 37.71 ± 0.60 15.92 ± 0.78 1.56 ± 0.72
36.65 ± 0.21 10.07 ± 0.13 2.59 ± 0.14 36.37 ± 0.2 14.08 ± 0.2 2.88 ± 0.13
5.04 ± 0.17 3.68 ± 0.17 3.84 ± 0.13 5.7 ± 0.17 2.53 ± 0.17 4.69 ± 0.13
13.17 ± 0.32 71.76 ± 0.4 2.37 ± 0.12 11.91 ± 0.32 46.57 ± 0.35 2.76 ± 0.09
– – – – 18.65 ± 0.21 –
6.11 ± 0.23 2.54 ± 0.11 89.72 ± 1.32 8.31 ± 0.21 2.24 ± 0.16 88.11 ± 1.34
B
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L. Lu et al. / Materials Characterization 117 (2016) 1–8
Fig. 3. Microstructures of steel A (a, c, e) and steel B (b, d, f) annealed for 1 h at: (a, b) 1000 °C; (c, d) 1050 °C; (e, f) 1100 °C.
less than that in Nb-free steels because Nb-alloying can decrease the contents of W, V while increase the contents of Fe, Cr in M2C carbides (Table 3).
Hardening is the most important heat treatment step for HSSs, and the austenitizing temperatures should be high enough to dissolve the carbides into the matrix, contributing to a high super saturation after
oil cooling. But the austenitizing temperature should ensure that the grain growth and the carbides aggregation would not be disastrous, thus the thermal stability is one of the important processing properties of HSSs. The grain size and the volume fraction of undissolved M6C carbides at different austenitizing temperatures are shown in Fig. 5. It shows the grain size is increased and the volume fraction of undissolved M6C carbides is decreased with increasing austenitizing temperature. However, the grain size of steel B increases slowly than that of steel A though
Fig. 4. XRD patterns of extracted carbides of steel A and B annealed at 1100 °C for 1 h.
Fig. 5. Variation of grain size and volume fraction of undissolved M6C carbides with austenitizing temperatures.
3.3. Effect of Nb on heat treatment response of SF HSSs
L. Lu et al. / Materials Characterization 117 (2016) 1–8
Fig. 6. Microsturctures of steel B (a, b, c, d) and steel A (e, f, g, h) hardening from temperatures of: (a, e) 1160 °C; (b, f) 1180 °C; (c, g) 1200 °C; (d, h) 1220 °C.
Fig. 7. Effect of tempering temperature on the mechanical properties of the investigated steels.
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Fig. 8. TEM micrographs and diffraction patterns of steel A (a, b, c) and steel B (d, e, f) tempered at 560 °C. Bright field (a, d); dark field (c, f).
the latter contains much more undissolved M6C carbides. For example, with increasing austenitizing temperature from 1180 °C to 1200 °C, the grain size of steel A is increased from 18.49 μm to 24.26 μm, for 31% increment, while that is 18% for steel B (from 18.23 μm to 21.54 μm). Fig. 6 illustrates the microstructures of steels hardening from different temperatures. Two types of carbide were distinguished by different contrast: white particals of irregular shape was M6C and the gray ones was MC. The dissolution, coagulation, polygonization of carbides during high-temperature austenitizing could be observed in Fig. 6. The rapid growth of grain (Fig.5) and serious aggregation of M6C (Fig. 6d and h) reveals that the austenitizing temperature of 1220 °C seems too high for both steel A and steel B, especially for steel A because the polygonization of carbides will increase the brittleness [34,35]. Compare Fig. 6d and h, it is obviously to find that the M6C aggregation degree is smaller in steel B. Base on the results of Figs. 5 and 6, a higher thermal stability of the steel B is evident. It is widely approve that Nb carbides have very poor solubility at hardening temperature [36], thus, Nb carbides can reduce the growth of the grain size and allow higher austenitization temperatures. 3.4. Mechanical properties Fig. 7 shows the effect of tempering temperature on the mechanical properties, including the bending strength, impact toughness and hardness. The bending strength of steel A follows a decreasing trend as the tempering temperature increases from 540 °C to 600 °C. While that of steel B raises at first, then decreases, and the peak value appears at 560 °C. The secondary hardening phenomenon of both two steels appears, with the maximum hardness value at 540 °C for steel A and
560 °C for steel B, respectively. Steel B has higher bending strength and hardness than steel A at this tempering temperature range except for 540 °C. Hardness as one of the important mechanical properties mainly depends on the secondary carbides precipitation during tempering. There is no evidence for Nb directly takes part in the secondary carbides precipitation. However, just as discussed previously, the compositions of MC carbides changed remarkably with the addition of Nb, about 25% of V is substituted by 18% of Nb, means that more vanadium dissolved in the matrix. On the other hand, the contents of W, Mo and Cr of the matrix are also higher in steel B (see Table 3). These alloying elements have an important contribution to the secondary hardness by precipitating finer secondary carbides [37,38]. Fig. 8 shows the secondary carbides and its diffraction patterns for steel A and B tempered at 560 °C. The very fine needle-shape secondary carbides with length of 8 ~ 12 nm and diameter b 3 nm, and with 2–3 nm average spacing were identified by electron diffraction as being M2C in Fig. 8a and c. As shown by the selected area diffraction pattern in Fig. 8b, the M2C carbides were related to the ferrite lattice by Pitsch-Schrader orientation relationship. The measured sizes of the needle-like carbides in steel B (Fig. 8d) are nearly as long as those in steel A, but with 1–2 nm average spacing, which shows that the dense of secondary carbides in steel A is less than that in steel B. Thus, it may conclude that the dispersion strengthening in steel B is stronger than that in steel A. It is also possible that relatively high concentration of dissolved alloying elements in the matrix of Nb-contained steel (Table 3) can contribute to the solid solution strengthening. All these possible factors contribute to the higher strength in steel B. The other important property of HSSs is toughness. As can be seen in Fig. 7, the impact toughness of steel B is lower than steel A at the same
Fig. 9. Fractographs of the investigate steels tempered at 560 °C. (a) Steel A; (b) steel B.
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ASP23 HSS. Thus, the spray-formed steel B has the best performance in these steels. 4. Conclusions (1) Spray forming and Nb-alloying can effectively refine the solidification microstructures (especially refining the M2C eutectic carbides) and reduce the amount of eutectic carbides of HSSs. Moreover, Nb addition can lower the stability of M2C carbides by decreasing the contents of W, Mo, V and increasing the content of Fe in M2C carbides. (2) Spray-formed Nb-alloyed high-speed steel has higher thermal stability, superior hardness and bending strength after tempering except for having slightly inferior impact toughness compared to that of Nb-free steel. (3) The tool made from spray-formed Nb-alloying HSS has the best performance than that made from conventional ingot casting and PM HSSs.
Fig. 10. Reamers made by HSSs.
Acknowledgements This work was financially supported by the National Basic Research Program of China (no. 2011CB606303).
tempering temperature except 600 °C. The impact toughness can be influenced by various factors [39,40]: (i) the kind and level of impurity; (ii) microstructural factors such as kind, size, volume fraction and distribution of carbides; and (iii) characteristics of the matrix. The fracture surfaces of both steels exhibit quasi-cleavage facets, tear ridges, and voids associated with undissolved carbides, as shown in Fig. 9. And the lower impact toughness of steel B could be attributed to the harder matrix and much higher volume fraction of hard primary MC carbides. (Fig. 9b).
3.5. Tool life evaluation To investigate the tool life made from spray formed HSSs, some reamers (Fig. 10) are machined from commercial AISI M2 HSS (made by conventional ingot cast), ASP23 HSS (made by powder metallurgy), and spray-formed HSSs with or without Nb addition. Fig. 11 shows the performances of the reamers in uninterrupted reaming, with almost the same hardness. It can be found that the tool life of spray-formed steel A (without Nb) is 1.8 times as long as that of commercial M2 HSS. The tool life of spray- formed steel B (with Nb) is ~ 3 times the length of the commercial AISI M2 HSS and 1.3 times as long as that of
Fig. 11. Tool lives of various HSSs.
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Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Improved the microstructures and properties of M3:2 high-speed steel by spray forming and niobium alloying L. Lu a, L.G. Hou a,⁎, J.X. Zhang a, H.B. Wang a, H. Cui a, J.F. Huang a, Y.A. Zhang b, J.S. Zhang a a b
State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Xueyuan Road 30, Haidian District, Beijing 100083, China State Key Laboratory of Non-Ferrous Metals and Process, General Research Institute for Non-Ferrous Metals, Beijing 100088, China
a r t i c l e
i n f o
Article history: Received 21 May 2015 Received in revised form 16 April 2016 Accepted 19 April 2016 Available online 20 April 2016 Keywords: High-speed steel Spray forming Mechanical properties Nb alloying
a b s t r a c t The microstructures and properties of spray formed (SF) high-speed steels (HSSs) with or without niobium (Nb) addition were studied. Particular emphasis was placed on the effect of Nb on the solidification microstructures, decomposition of M2C carbides, thermal stability and mechanical properties. The results show that spray forming can refine the cell size of eutectic carbides due to the rapid cooling effect during atomization. With Nb addition, further refinement of the eutectic carbides and primary austenite grains are obtained. Moreover, the Nb addition can accelerate the decomposition of M2C carbides and increase the thermal stability of high-speed steel, and also can improve the hardness and bending strength with slightly decrease the impact toughness. The high-speed steel made by spray forming and Nb alloying can give a better tool performance compared with powder metallurgy M3:2 and commercial AISI M2 high-speed steels. © 2016 Elsevier Inc. All rights reserved.
1. Introduction High-speed steels (HSSs), as a special class of highly alloyed tool steels, are extensively used as cutting tools, cold molds and rolls for the exceptional combination of excellent hot hardness, high wear resistance, well toughness and good workability. These properties are mainly depended on the size, type, distribution, shape and amount of carbides and the state of the matrix [1–3]. HSSs made by conventional ingot casting contain coarse network of eutectic carbides formed at grain boundaries. Even after certain forging degree, the carbide particles are arranged in bands parallel to forging direction, which deteriorating the isotropic properties and toughness of HSSs [4,5]. On the other hand, the slow cooling rate of conventional ingot casting limited the addition of strong carbide forming elements such as Ti, Nb, V due to the formation of coarse carbides. Fortunately, the rapid solidification technology provides a new opportunity to modify the solidification microstructures and improve the properties of HSSs, e.g., powder metallurgy (PM) was the first and the most mature industrial application of the rapid solidification technology for highspeed steel products. Finer and more uniform primary carbides, smaller grain sizes and the absence of carbides stringers and macro-segregation are some characteristics attained by PM HSSs compare with the conventional ingot casting HSSs [6–8]. Also it is possible to produce any alloy design by PM process. However, the PM HSSs is still applied in a narrow field because they have normally high cost due to the complex and rigorous processing steps. ⁎ Corresponding author at: University of Science and Technology Beijing, Beijing, China. E-mail address: [email protected] (L.G. Hou).
http://dx.doi.org/10.1016/j.matchar.2016.04.010 1044-5803/© 2016 Elsevier Inc. All rights reserved.
Spray forming (SF), a promising preparation technology for materials, can combine the rapid solidification (gas atomization) and nearnet-shape forming (deposition) and has been proved to be an effective method to improve the solidification microstructures of tool steels [9–12], due to the high cooling rate (103–105 K/s) during gas atomization and the microstructural deformation or fracture in deposition stage [13]. In contrast to conventional ingot casting, the most significant microstructure characteristics of SF HSSs are the smaller equiaxed grains, finer and much more uniform-distributed carbides, and less degree of segregation [6,8], resulting in the improved performance of SF HSSs. However, the increase of cooling rate in atomization stage is limited by the density of the final billet [14–16]. Alloying design has been widely used to improve the microstructures and properties of HSSs, and two main alloying strategies are proposed: adding non-carbide forming elements (e.g., Co, Al, Si and rare earth) [17–19], or modifying the microstructures and properties with strong carbide forming elements (e.g., Ti, Nb, Zr). Recently, a relatively great number of literatures reflect the enhanced interest of researchers in various countries on the use of Nb and Ti for alloying HSSs [20–27], but most of the mentioned works were based on conventional ingot casting. In this paper, the idea of combining spray forming and Nb alloying to improve the performance of M3:2 HSS is implemented and the purpose of the present work is to study the microstructures and mechanical properties of spray formed Nb-containing HSSs as well as the role of Nb in HSSs. 2. Experimental procedure Billet with dimension of Φ180 mm × 100 mm was prepared by spray forming and the processing parameters are shown in Table 1. A
2
L. Lu et al. / Materials Characterization 117 (2016) 1–8 Table 1 Process parameters of spray forming. Items
Parameters
Super heat/°C Diameter of the delivery tube/mm Atomization gas Atomization pressure/MPa Deposition distance/mm
160–180 4.0 N2 0.45–0.5 425–450
Table 2 Chemical compositions of the investigated steels (wt.%).
Steel A Steel B
C
W
Mo
Cr
V
Nb
1.30 1.31
6.20 6.10
5.10 4.90
4.60 4.48
2.80 2.75
– 0.50
commercial AISI M2 HSS used as the feedstock was melted with appropriate additions of alloying elements. The chemical compositions of the experimental steels are listed in Table 2. A cylindrical specimen with a size of Φ50 mm × 40 mm was cut from the center of the spray-formed billet and subjected to isothermal annealing at 1180 °C for 1.5 h, then forged to Φ16 mm bars at 1150 °C followed with slowly cooling to room temperature. Before the hardening treatment, all the specimens were heated to 900 °C with a heating rate of 100 °C/h and soaked for 1.5 h, then cooled to 25 °C with a cooling rate of 20 °C/h. Subsequently, the specimens were austenitized at 1160–1220 °C for 25 min and oil quenched to room temperature, followed by triple tempering at 540–600 °C for 1 h. In order to prevent decarburization during heat treatment, each sample was sealed in vacuum quartz tube. The microstructures were characterized using field emission scanning electron microscope (SEM), and Energy dispersive X-ray spectroscopy (EDS) was employed for local chemical composition determination and for the distinction of MC and M2C carbides. For metallographic examinations, the samples were prepared by grinding and polishing, and were etched with 8% nital solution. The carbide volume fraction as well as grain size was measured by using image analysis software (Image-pro plus 6). The reported values for each sample were
taken from at least 50 measurements in different fields with a magnification of 1000 times. The extraction of carbides was performed using electrolysis, operating at 40 V, 0 °C with a electrolyte contained 7 g citric acid, 20 mL hydrochloric acid and 250 mL methanol. X-ray scattering was made with a Rigaku X-ray Diffractometer (D/MAX-RB) with CuKα radiation (λ = 1.5406 Å) at 12 kW. Differential scanning calorimetry (DSC) experiments were carried out in a Netzsch DSC calorimeter with cooling rate of 10 K/min from 1455 to 200 °C. The DSC samples (Φ5 mm × 1 mm) were machined from the asdeposited billet. Three-point bending tests were carried out on specimens with dimension of 5 × 5 × 35 mm at a bending rate of 0.1 mm/min. The sample size for impact toughness tests was 10 × 10 × 55 mm without notch according to ISO 5754. The bulk hardness measurements were carried out using Rockwell “C” scale with a load of 150 kgf (589 N) applied for 3 s and 7–10 readings were considered for estimating the hardness values. The reamers machined from the forged bars were sealed in vacuum quartz tube and annealed at 900 °C for 1.5 h, and then austenitized at 1200 °C for 25 min and oil-quenched to room temperature, followed by triple tempering at 560 °C for 1 h. The reamers used to ream the medium carbon steel with average hardness of 32HRC, and all the tests were under the condition of oil lubrication. 3. Results and discussion 3.1. Effects of spray forming and Nb alloying on the microstructures of HSSs The microstructures of the conventional ingot casting HSS and SF HSS in Fig. 1 show that the solidification microstructures are consisted of equiaxed grains, discontinuous network of plate-shaped M2C eutectic carbides (white) and MC carbides (gray) both at grain boundaries and within the matrix. It can be obtained that the grain size of SF HSS is similar to that of conventional ingot casting HSS (~60 μm). However, the sizes of M2C and MC carbides in SF HSS are remarkably refined (Fig. 1a and b), e.g., the cell sizes of M2C eutectic carbides in conventional ingot casting HSS and SF HSS are ~35 μm, ~17 μm, respectively. R.A. Mesquita and C.A. Barbosa [4] reported that high cooling rate of atomization is the main factor for primary carbides refining. Thus, the cooling rate could be the key factor to refine the solidification microstructure of
Fig. 1. Microstructures of steel A (a, b) and B (c) made by ingot cast (a) and spray forming (b, c).
L. Lu et al. / Materials Characterization 117 (2016) 1–8
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particles in steel B can be act as the inoculants for the nucleation during eutectic reaction and refine the eutectic carbides. These two factors are the main reason for the refinement of M2C carbides. From the composition of MC and M2C carbides in the spray-formed HSSs (Table 3), it is can be found that Nb alloying can decrease the contents of W, V while increase the content of Fe in M2C carbides, which will be beneficial to the decomposition of M2C carbides (in Section 3.2). Also the composition of MC carbides is changed remarkably with Nb addition, about 25% of V is substituted by 18% of Nb, and the content of W and Mo are increased. 3.2. Effect of Nb on the decomposition of M2C carbides
SF HSSs. But much higher cooling rates may decrease the liquid fraction of atomized particles and eventually result in the density reduction of the as-sprayed billet. In order to further refine the as-deposited microstructures, Nb is added and Fig. 1c shows excellent microstructures which can be characterized as fine grains, uniform distribution of fine M2C eutectic carbides and MC carbides. The mean cell sizes of M2C eutectic carbides in Fig. 1c is ~9 μm and the grain size is 45 μm. The volume fraction of M2C eutectic carbides decreases from 5.33% to 4.05% and the volume fraction of MC carbides increases from 2.68% to 3.12% with Nb addition (Fig. 1b and c). Generally, it can be concluded that the spray-formed HSSs can have much finer microstructures and less degree of carbides segregation than that of conventional ingot casting HSSs. Adding Nb also can considerably refine the prior-austenite grain sizes, reduce the cell size and volume fractions of M2C eutectic carbides while increase the amounts of MC carbides. Thus, the spray forming process combining with Nb alloying can effectively modify the solidified microstructures of HSSs. Basing on the results of ref [28,29], the solidification sequences of M3:2 high-speed steel with high carbon content can be described as follows:
M2C eutectic carbide which is a meta-stable phase in as-deposited microstructures will deteriorate the hot workability and toughness [5]. According to practical experience we expect that it has been decomposed prior to hot working. There are some reports [30–33] indicate that M2C carbides will decompose into M6C and MC carbides when annealed at high temperatures (N1000 °C). It is conceivable that using the decomposition treatment to obtain a more uniform carbides size and distribution, indeed to improve the performance of HSSs. To explore this possibility, it is necessary to study the effect of Nb on the decomposition of M2C carbides. Fig. 3 shows the microstructures of spray formed steels annealed at different temperatures for 1 h. At 1000 °C, no decomposition of M2C carbides can be found in steel A (Fig. 3a) and steel B (Fig. 3b). Increasing the annealing temperature to 1050 °C, most of the M2C carbides in steel A are still not decomposed except a few of small sized M2C carbides distributed on the edge of eutectic cells (arrowed in Fig. 3c). While all of the M2C carbides in steel B are decomposed into M6C and MC carbides, and the MC carbides are distributed as small dots inside the M6C carbides (Fig. 3d). But the new carbides still keep the original plate-shape of M2C carbides. At 1100 °C, partial decomposition of large sized M2C carbides in steel A (Fig. 3e) are obtained. But in steel B, the separation occurs at the interface between M6C and MC carbides, and the necking areas are formed in the thinner regions of the plates (Fig. 3f). The XRD results of the extracted carbides (Fig. 4) also show the peaks of M2C carbides in steel A, no peaks of M2C carbides were detected in steel B, meaning the complete decomposition of M2C carbides at 1100 °C in steel B. The decomposition reaction of M2C carbides obeys the formula as follows [31]:
1. Precipitation of austenite from liquid.
M2 C þ γ→M 6 C þ MC:
Fig. 2. DSC curves for investigated steels cooled from 1450 °C.
2. Formation of carbides at the interface of γ and liquid. 3. Formation of eutectic carbides through eutectic reaction.
ð1Þ
H. Fredriksson [33] observed that the nucleation of M6C carbides occurred firstly at M2C/austenite interfaces, and the transformation would be proceeded with the nucleation of MC carbides at either M6C/austenite or M2C/M6C interfaces, and the decomposition rate was proposed to depend upon the diffusion of alloying elements. Mainly, M6C carbides need Fe atoms from the austenite and Mo/W atoms from M2C. At a constant decomposition treating temperature and time, the finer M2C plates will be decomposed quickly than the coarse ones. As mentioned in Section 3.1, Nb alloying can refine the M2C carbides, so the decomposition of M2C carbides in steel B could be faster than that in steel A. Moreover, the stability of M2C carbides in Nb-alloying steels will be
The DSC curves in Fig. 2 show that the presence of 0.5% Nb in M3:2 HSS will not change the solidification sequences, but it apparently increases the temperature for MC formation and decreases the liquidus temperature. That is to say the primary austenite precipitates at a lower temperature in steel B, thus the Nb containing steel has a finer austenite grains than Nb free steel (see Fig. 1b and c). Because of high affinity with carbon, Nb reacts with free carbon of the melt, decreasing the carbon content of the melt. In such a case, the amount of residual interdendritic liquid for eutectic reaction will decrease, and more MC
Table 3 Composition of primary carbides and the matrix (at.%). Steel
Carbides
W
Mo
Cr
V
Nb
Fe
A
M2C MC Matrix M2C MC Matrix
39.02 ± 0.54 11.94 ± 0.61 1.48 ± 0.69 37.71 ± 0.60 15.92 ± 0.78 1.56 ± 0.72
36.65 ± 0.21 10.07 ± 0.13 2.59 ± 0.14 36.37 ± 0.2 14.08 ± 0.2 2.88 ± 0.13
5.04 ± 0.17 3.68 ± 0.17 3.84 ± 0.13 5.7 ± 0.17 2.53 ± 0.17 4.69 ± 0.13
13.17 ± 0.32 71.76 ± 0.4 2.37 ± 0.12 11.91 ± 0.32 46.57 ± 0.35 2.76 ± 0.09
– – – – 18.65 ± 0.21 –
6.11 ± 0.23 2.54 ± 0.11 89.72 ± 1.32 8.31 ± 0.21 2.24 ± 0.16 88.11 ± 1.34
B
4
L. Lu et al. / Materials Characterization 117 (2016) 1–8
Fig. 3. Microstructures of steel A (a, c, e) and steel B (b, d, f) annealed for 1 h at: (a, b) 1000 °C; (c, d) 1050 °C; (e, f) 1100 °C.
less than that in Nb-free steels because Nb-alloying can decrease the contents of W, V while increase the contents of Fe, Cr in M2C carbides (Table 3).
Hardening is the most important heat treatment step for HSSs, and the austenitizing temperatures should be high enough to dissolve the carbides into the matrix, contributing to a high super saturation after
oil cooling. But the austenitizing temperature should ensure that the grain growth and the carbides aggregation would not be disastrous, thus the thermal stability is one of the important processing properties of HSSs. The grain size and the volume fraction of undissolved M6C carbides at different austenitizing temperatures are shown in Fig. 5. It shows the grain size is increased and the volume fraction of undissolved M6C carbides is decreased with increasing austenitizing temperature. However, the grain size of steel B increases slowly than that of steel A though
Fig. 4. XRD patterns of extracted carbides of steel A and B annealed at 1100 °C for 1 h.
Fig. 5. Variation of grain size and volume fraction of undissolved M6C carbides with austenitizing temperatures.
3.3. Effect of Nb on heat treatment response of SF HSSs
L. Lu et al. / Materials Characterization 117 (2016) 1–8
Fig. 6. Microsturctures of steel B (a, b, c, d) and steel A (e, f, g, h) hardening from temperatures of: (a, e) 1160 °C; (b, f) 1180 °C; (c, g) 1200 °C; (d, h) 1220 °C.
Fig. 7. Effect of tempering temperature on the mechanical properties of the investigated steels.
5
6
L. Lu et al. / Materials Characterization 117 (2016) 1–8
Fig. 8. TEM micrographs and diffraction patterns of steel A (a, b, c) and steel B (d, e, f) tempered at 560 °C. Bright field (a, d); dark field (c, f).
the latter contains much more undissolved M6C carbides. For example, with increasing austenitizing temperature from 1180 °C to 1200 °C, the grain size of steel A is increased from 18.49 μm to 24.26 μm, for 31% increment, while that is 18% for steel B (from 18.23 μm to 21.54 μm). Fig. 6 illustrates the microstructures of steels hardening from different temperatures. Two types of carbide were distinguished by different contrast: white particals of irregular shape was M6C and the gray ones was MC. The dissolution, coagulation, polygonization of carbides during high-temperature austenitizing could be observed in Fig. 6. The rapid growth of grain (Fig.5) and serious aggregation of M6C (Fig. 6d and h) reveals that the austenitizing temperature of 1220 °C seems too high for both steel A and steel B, especially for steel A because the polygonization of carbides will increase the brittleness [34,35]. Compare Fig. 6d and h, it is obviously to find that the M6C aggregation degree is smaller in steel B. Base on the results of Figs. 5 and 6, a higher thermal stability of the steel B is evident. It is widely approve that Nb carbides have very poor solubility at hardening temperature [36], thus, Nb carbides can reduce the growth of the grain size and allow higher austenitization temperatures. 3.4. Mechanical properties Fig. 7 shows the effect of tempering temperature on the mechanical properties, including the bending strength, impact toughness and hardness. The bending strength of steel A follows a decreasing trend as the tempering temperature increases from 540 °C to 600 °C. While that of steel B raises at first, then decreases, and the peak value appears at 560 °C. The secondary hardening phenomenon of both two steels appears, with the maximum hardness value at 540 °C for steel A and
560 °C for steel B, respectively. Steel B has higher bending strength and hardness than steel A at this tempering temperature range except for 540 °C. Hardness as one of the important mechanical properties mainly depends on the secondary carbides precipitation during tempering. There is no evidence for Nb directly takes part in the secondary carbides precipitation. However, just as discussed previously, the compositions of MC carbides changed remarkably with the addition of Nb, about 25% of V is substituted by 18% of Nb, means that more vanadium dissolved in the matrix. On the other hand, the contents of W, Mo and Cr of the matrix are also higher in steel B (see Table 3). These alloying elements have an important contribution to the secondary hardness by precipitating finer secondary carbides [37,38]. Fig. 8 shows the secondary carbides and its diffraction patterns for steel A and B tempered at 560 °C. The very fine needle-shape secondary carbides with length of 8 ~ 12 nm and diameter b 3 nm, and with 2–3 nm average spacing were identified by electron diffraction as being M2C in Fig. 8a and c. As shown by the selected area diffraction pattern in Fig. 8b, the M2C carbides were related to the ferrite lattice by Pitsch-Schrader orientation relationship. The measured sizes of the needle-like carbides in steel B (Fig. 8d) are nearly as long as those in steel A, but with 1–2 nm average spacing, which shows that the dense of secondary carbides in steel A is less than that in steel B. Thus, it may conclude that the dispersion strengthening in steel B is stronger than that in steel A. It is also possible that relatively high concentration of dissolved alloying elements in the matrix of Nb-contained steel (Table 3) can contribute to the solid solution strengthening. All these possible factors contribute to the higher strength in steel B. The other important property of HSSs is toughness. As can be seen in Fig. 7, the impact toughness of steel B is lower than steel A at the same
Fig. 9. Fractographs of the investigate steels tempered at 560 °C. (a) Steel A; (b) steel B.
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ASP23 HSS. Thus, the spray-formed steel B has the best performance in these steels. 4. Conclusions (1) Spray forming and Nb-alloying can effectively refine the solidification microstructures (especially refining the M2C eutectic carbides) and reduce the amount of eutectic carbides of HSSs. Moreover, Nb addition can lower the stability of M2C carbides by decreasing the contents of W, Mo, V and increasing the content of Fe in M2C carbides. (2) Spray-formed Nb-alloyed high-speed steel has higher thermal stability, superior hardness and bending strength after tempering except for having slightly inferior impact toughness compared to that of Nb-free steel. (3) The tool made from spray-formed Nb-alloying HSS has the best performance than that made from conventional ingot casting and PM HSSs.
Fig. 10. Reamers made by HSSs.
Acknowledgements This work was financially supported by the National Basic Research Program of China (no. 2011CB606303).
tempering temperature except 600 °C. The impact toughness can be influenced by various factors [39,40]: (i) the kind and level of impurity; (ii) microstructural factors such as kind, size, volume fraction and distribution of carbides; and (iii) characteristics of the matrix. The fracture surfaces of both steels exhibit quasi-cleavage facets, tear ridges, and voids associated with undissolved carbides, as shown in Fig. 9. And the lower impact toughness of steel B could be attributed to the harder matrix and much higher volume fraction of hard primary MC carbides. (Fig. 9b).
3.5. Tool life evaluation To investigate the tool life made from spray formed HSSs, some reamers (Fig. 10) are machined from commercial AISI M2 HSS (made by conventional ingot cast), ASP23 HSS (made by powder metallurgy), and spray-formed HSSs with or without Nb addition. Fig. 11 shows the performances of the reamers in uninterrupted reaming, with almost the same hardness. It can be found that the tool life of spray-formed steel A (without Nb) is 1.8 times as long as that of commercial M2 HSS. The tool life of spray- formed steel B (with Nb) is ~ 3 times the length of the commercial AISI M2 HSS and 1.3 times as long as that of
Fig. 11. Tool lives of various HSSs.
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