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VN composite powders
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Dense M2 high speed steel containing core-shell MC carbonitrides using high-energy ball milled M2/VN composite powders Nan Chen, Ren Luo, Huiwen Xiong **, Zhiyou Li * State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, China
A R T I C L E I N F O
A B S T R A C T
Keywords: High speed steel Microstructure High-energy ball milling Sintering Carbonitride
Uniform M2/VN composite powders were prepared via high-energy ball milling of commercial M2 high-speed steel and VN powders, and a promising M2/VN high speed steel with 4 wt% V was obtained by vacuum sin tering. The phases, microstructure of the milled composite powders, and densification behavior, related me chanical properties of sintered steels were investigated. The VN particles were embedded in the matrix phase α-Fe during the ball milling process. After 28 h of ball milling, M6C carbides were continuously refined and showed a streamline distribution. Nearly full densification (~99.4% RD) of M2/VN steel was achieved at 1160 � C by supersolidus liquid phase sintering. Carbonitrides with core-shell structure, exhibiting high N content, were confirmed via the dissolution and precipitation process. The VN particles effectively inhibited the grain growth and improved both the hardness and bend strength. Attracting M2 high speed steel with fine and uniformly dispersed MC carbides was achieved, showing satisfactory bending strength of 3000 MPa and hardness of 62.2 HRC.
1. Introduction 1.1. PM HSS based on the SLPS mechanism High-speed steel (HSS) has been widely used as cutting tools, high load dies, wear-resistant parts due to its excellent mechanical properties [1]. Compared with the conventional cast or wrought parts, HSS fabri cated using powder metallurgy technology (PM HSS) exhibits improved wear resistance, hardness, heat resistance and toughness, owing to the finer and more homogeneous microstructure. In most cases, the densi fication of the sintered PM HSS is achieved through a super-solidus liquid-phase sintering (SLPS) mechanism [2]. The optimum sintering temperature is between the temperature range of the austenite, carbide phase and liquid phase. Low sintering temperature leads to insufficient liquid phase and unsatisfactory density, while high temperature results in the formation of coarser grains and deteriorated performance [3]. General, the main issues of SLPS are as follows: 1) The presence of liquid phase promotes the precipitation of unsatisfactory carbides with mor phologies of rod-like or coarser along the grain boundary [4]; 2) The sintering temperature is relatively high with a narrow window. 3) Careful controlling of the sintering process is needed for achieving the
desired phase components with fine structure [5]. 1.2. M2 HSS with VN particles via vacuum sintering V-containing PM HSS belongs to an attracting material, owing to the in-situ formed fine MC carbide and related high hardness, toughness and wear resistance [6]. As previously reported [7–9], in the nitrogen-rich atmosphere, the sintering temperature of PM HSS with high V content (over 10 wt%) decreases. The N element also results in the enlarged sintering window and the formation of fine carbonitrides and nitrides. The T42 HSS sintered in the nitrogen-rich atmosphere decreased the optimum sintering temperature (OST) from 1220 � C (that of vacuum sintering) to 1160 � C, and widened the sintering window (SW) nearly 30–40 � C [10]. Hence, the preparation of PM HSS with high V content via sintering in nitrogen atmosphere shows a way to handle the above issues of SLPS. Nevertheless, for PM HSS with relatively low V element content, the nitrogen atmosphere sintering is not applicable, owing to the decreased sinterability [7]. The OST of low vanadium high speed steels of M2 HSS (around 2.0 wt% V element content) increased from 1110 � C to 1230 � C when sintered under the nitrogen-rich atmosphere [7]. Based on this, herein, nitrogen was introduced by directly adding
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (H. Xiong), [email protected] (Z. Li). https://doi.org/10.1016/j.msea.2019.138628 Received 12 September 2019; Received in revised form 1 November 2019; Accepted 2 November 2019 Available online 4 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: Nan Chen, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138628
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Fig. 1. Microscopic morphology of the M2 HSS powders (a) and VN powders (b). Table 1 Chemical composition of the M2 HSS powders (mass fraction, %). C
W
Mo
Cr
V
Fe
0.8
6.6
5.2
4.1
1.9
bal.
Table 2 Chemical composition of the VN powders (mass fraction, %). C
V
N
O
3.0
75.6
19.8
bal.
VN particles into the HSS powders, and M2/VN-based HSS with a low V element content of 4 wt% was fabricated via vacuum sintering. The research scheme here may be a promising way to obtain fine MC rein forced PM HSS, and avoid the problem of poor sinterability induced by the nitrogen-rich atmosphere.
Fig. 2. X-ray diffraction patterns of composite powders after milling.
and enhance the mechanical properties of materials [24]. However, there are few reports on the direct sintering preparation process of high-energy ball-milled water atomized HSS powders. The sinterability of M2 HSS is notoriously difficult, and direct sin tering route seems not suitable for it from the past work [25]. Significant improvement of its sinterability is unlikely to be successful through minor modification to its composition or a change to the sintering at mosphere and forming method [26,27]. This paper describes an inves tigation into improving the sinterability of M2 HSS powder by mechanical alloying and to enhance mechanical properties by rein forcement with VN particles. In this research, VN/M2 composite pow ders were prepared using high-energy ball milling. Microstructure evolution of composite powders during the milling process was studied. The densification behavior and mechanical properties of samples are investigated.
1.3. High energy ball milling for PM HSS containing MC Many studies [11–15] have shown that properties of PM HSS were further improved by increasing the amount of carbide forming elements to increase the volume fraction of the hard phases, or by directly adding stable hard phase particles. For instance, the hardness and wear resis tance of M2 HSS are improved by adding carbide particles such as VC [12–14], TiC [16] and NbC [17]. The stability and hardness of MC-type carbide in HSS are higher than those of other carbides such as M6C-type carbide [18], so increasing the content of MC-type carbide is more ad vantageous to improve the hardness and red hardness of materials. When carbide particles are mixed into HSS powder mechanically, the hard phase tends to aggregate between the steel powders, which not only reduces the sintering strength of the materials, but also produces coarse carbides in SLPS [19]. High-energy ball milling has advantages in introducing reinforcing phase, which can refine the powders and disperse the additives evenly into the powder particles [16]. High-energy ball milling generally refers to the process of reducing powder size by mechanical energy [20]. This technology has been suc cessfully applied to synthesize a lot of commercial alloys and composite powders [21–23]. Liu et al. [16]. Refined the powder of M2 HSS by high-energy ball milling and realized the fine dispersion distribution of TiC hard particles in M2 HSS. High energy ball mill can refine the grain structure and make the powders have higher lattice distortion energy. Thus, high-energy ball milling can improve the sinterability of materials
2. Experimental materials and methods 2.1. High-energy ball milling of composite powders The SEM morphology of M2 and VN powders is shown in Fig. 1. The irregular shaped powders with mean particle size of 75 μm were water atomized M2 high speed steel powders, which were provided by €gana €s China Co., Ltd. The VN powders particles are polygonal in Ho shape and they were purchased from Shanghai Yunfu Nanotechnology 2
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Fig. 3. SE SEM images of composite powders milled for (a) 2 h; (b) 4 h; (c) 8 h; (d) 16 h; (e), (f) 28 h.
Co., Ltd. The chemical composition of the M2 HSS powders and VN powders is given in Table 1 and Table 2, respectively. The two kinds of powders were formulated into composite powders with the vanadium element content of 4 wt%. The VN/M2 composite powders were milled by a planetary ball mill with a ball-to-powder weight ratio of 20:1 and a rotating speed of 300 rpm. The ball mill bowls and milling balls are made of stainless steel and GCrl5 alloy, respectively. Argon was used as the protective atmo sphere, and the ball milling time was kept at 2, 4, 8, 16 and 28 h. After the high-energy ball milling, 4.0 wt% of polyethylene glycol, 2.0 wt% of
colloidal graphite and 50 ml of ethanol were added to the composite powders. Polyethylene glycol was dissolved in ethanol before adding into the composite powders. Subsequently, the mixtures were mixed at 150 rpm for 2 h to ensure the adequate homogenization. After ballmilling, drying of the composite powders was carried out in a vacuum oven at 80 � C for 8 h. Finally, the dried composite powders were passed through a 200-mesh sieve for further handling.
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Fig. 4. Internal microstructures of composite powders milled for (a) 2 h; (b) 4 h; (c) 8 h; (d) 16 h; (e), (f) 28 h.
2.2. Preparation of the samples
2.3. Analysis and measurements
The prepared composite powders mixtures were cold-pressed at a pressure of 200 MPa with a dimension of 25 mm � 8 mm � 8 mm. Then, the prepared specimens were put in a tubular hydrogen furnace under flowing H2 gas, and heated at 400 � C for 60 min to remove the binder. The debound samples were placed into the vacuum sintering furnace and first sintered at 1030 � C for 1 h. The purpose of the step above is deoxidize the samples completely. Finally, the sintering of samples was carried out at 1150, 1160, 1180, 1200 and 1220 � C for 1 h, and then cooled to room temperature in the furnace. The heating rate of all the above heating processes was 5 � C/min.
The composite powders and the sintered samples were subjected to phase analysis using a D/Ma2500VB-RA X-ray diffractometer. A FEI Quanta 250 FEG field emission scanning electron microscope (SEM) and a PHOENIX energy dispersive spectrometer (EDS) were used to observe the microstructure and analyze the different phases. Elemental analysis of samples was performed by a JXA-8230 field emission electron probe microanalyzer (EPMA). The grain orientation and size of samples were analyzed using a FEI Helios Nanolab G3 UC electron backscattered diffraction (EBSD) system. A LE COCS-444C/S analyzer and a LECO TC436 N/O analyzer were used to measure carbon and oxygen contents, respectively. Densities of samples were measured by the Archimedes’ 4
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Fig. 5. Internal micro-area N element plane scan of the composite powders. Milled for 28 h; (a) BSED-SEM image of the composite powders; (b) N element dis tribution map.
approximate with each other, their diffraction peaks overlap in Fig. 2. It can be seen from Fig. 3(a) that the shape of VN powders was un changed and the water atomized M2 HSS powder particles were squashed from the original irregular shape after 2 h ball milling. After 4 h of ball milling, the M2 powders became flat and the edges tend to be smooth since the powders were soft at this stage. After 8 h, the powders were obviously broken and refined into fragments, and the degree of edge roundness increased. After 16 h, the powders were further refined, and the number of small particles increased significantly. After 28 h of milling, a large number of fine particles were formed, and the particles tended to be spherical. At the same time, there were aggregates of fine particles. It can be seen from Fig. 3(e) that a large number of fine par ticles were flakily welded on the surface of the large particles, and a depression caused by the peeling of fine particles from the surface of the large particles was observed, indicating that the powders have a mutual balance between crushing and cold welding during ball milling [28]. In order to reveal the phases and morphology of powders, the in ternal microstructures of composite powders are shown in Fig. 4. It can be seen that the initial M2 powder contains white granular M6C carbides and dark grey MC carbides, and most of MC carbides are adjacent to M6C carbides. After 4 h of ball milling, the dark grey block VN was embedded into the M2 matrix, and the MC carbides were separated from the M6C carbides. After 16 h, cracks and pores appeared inside the powder, which may be caused by the crushing and falling off of the hard and brittle phase particles during the milling. In the process of milling for 16 to 28 h, M6C carbides were continuously refined and tended to be lamellar, showing a streamline distribution, and its content was signif icantly reduced. The MC carbides were refined at the initial stage of ball milling, and the change of size and content during 8 h and subsequent ball milling was not obvious, indicating that MC carbides were stable during ball milling [6]. Fig. 5 shows the N-element mapping results of the internal micro-area of the powders after ball milling for 28 h. The N element was dispersedly distributed in M2 HSS, indicating that it was advantageous to prepare a uniform N-containing M2 HSS through the high energy ball milling.
Fig. 6. X-ray diffraction patterns of the sintered samples.
method. Rockwell hardness was provided with a load of 150 kg by a 500MRA-electric hardness tester. Bend strength was measured by three point bending with a span of 14.5 mm. Each numerical value was repeated at least three times and the results were expressed using the mean value. 3. Results and discussion 3.1. Phases, and microstructure of the M2/VN powders after milling Fig. 2 shows the X-ray diffraction patterns of the composite powders with different milling times. The initial powder mixtures mainly include α-Fe, M6C and M (C, N). As the ball milling time increased, the diffraction peaks were significantly broadened and weakened due to the refined crystal lattice and internal stress. In the middle and late stages of ball milling, some of the M6C and M (C, N) peaks disappeared. It is noteworthy that the change of diffraction peaks of carbides is not obvious because of their small volume fraction. The enlarged picture of the diffraction peaks between 42� and 45� is shown in Fig. 2. The first strong peak of Fe gradually shifted to the left, indicating that W, Mo and other large atoms of M6C were continuously dissolved into α-Fe. Since VC and VN are both cubic structures and the cell parameters are
3.2. Microstructure evolution during sintering XRD patterns of sintered samples are presented in Fig. 6. The M2/VN alloyed samples show obvious diffraction peaks of M6C and M (C, N) after sintering at 1150 � C. The reason is that the lattice distortion is eliminated after sintering and the diffraction peak becomes sharper [29]. Compared with the sample sintered at 1150 � C, the diffraction peak of Fe is obviously weakened and broadened after sintering at 5
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Fig. 7. BSED-SEM images of samples sintered in vacuum atmosphere at (a), (b) 1150 � C; (c), (d) 1160 � C; (e) 1180 � C; (f) 1200 � C.
1160 � C and above. While the peaks of M6C become weaker, the peaks of MC become stronger, and some peaks of M6C disappear, indicating that the content of M6C decreases and the content of MC increases. Mean while, the diffraction peak of α-Fe shifted to the higher angles. This is because the interplanar distance decreased due to dissolution of large atomic radius alloy elements in the matrix [30]. The microstructure evolution of the samples sintered at different temperatures is shown in Fig. 7. When the samples were sintered at 1150 � C, Closed circular pores remained at the boundary between the carbides and matrix as shown in Fig. 7(a) and (b). The size of M6C carbides (white phase) was significantly larger than that of M (C, N) and MC (grey phase). In most cases, High speed steel densifies via
supersolidus liquid phase sintering (SLPS). For supersolidus liquid phase sintering, sufficient liquid is required to enhance the densification of samples [16]. However, due to the gravity of samples, excess liquid can cause loss of slump and dimensional control. For supersolidus liquid phase sintering of high-speed steel, usually the amount of liquid as well as the densification rate and microstructure changes is controlled by the sintering temperature typically. The liquid phase has greater chance of appearing firstly at the interfaces between the enhanced particles and the matrix [16]. Carbide reinforcements and γ-Fe matrix form liquid phase through the reverse of eutectic reaction at the interfaces (γ-Fe þ MxCy→L) [31]. Some liquid phase usually migrates along the surface of γ-Fe matrix particles to the sintered neck, while the γ-Fe matrix that does 6
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shown in Fig. 7(d), the grey phases corresponded to a vanadium-rich carbonitride formed as core-shell structures with the VN phases as core and MC carbides as shell [32]. The shell formed through the dissolution and precipitation of original carbides in the M2 HSS [31]. When the temperature rose to 1180 � C, the carbides grew up signifi cantly. Many closed circular pores formed at the interfaces between carbides and the matrix. In addition, the carbides with core-shell structure have largely disappeared. At the temperature of 1200 � C, abnormal growth of MC carbide was observed. More liquid phase accelerated the rate of atomic diffusion. Through the dissolution-precipitation mechanism, smaller carbide dissolved in the liquid phase and then precipitated on the surface of larger carbide. The appearance of crescent carbides was the evidence of massive liquid phase in sintering process [33]. The graph of the average M6C carbide size as a function of temperature is shown in Fig. 8. The size of M6C carbide continued to increase from 0.89 μm to 1.38 μm along with the rising of the temperature. Fig. 9 shows the EPMA images of the sample sintered at 1160 � C. M (C, N) phase has high vanadium content in the central region and low vanadium content in the surrounding area. In the central region of M (C, N) phase, there are nitrogen-rich regions which are in red color. In the grey M (C, N) carbide, there is a gradient distribution of elements, which is corresponding to the reported non-uniform distribution of the MC phase in fully alloyed HSS prepared by gas atomization-hot isostatic pressing [30]. This also indicates that carbonitrides having a core of VN formed during the sintering process. Fig. 10 shows the EBSD grain distribution image and inverse pole figure (IPF) orientation color scheme of the sample sintered at 1160 � C. Black regions are caused by long-term vibration polishing resulting in carbide bulge. The M6C carbide is mostly clumpy and the MC carbide is mostly near-spherical as shown in Figs. 10 and 7(d). This reason may be that the M6C carbide directly precipitates from the liquid phase, while
Fig. 8. Average M6C carbide size of the samples at different temperatures.
not react with carbide reinforcements is still solid phase sintering. Thus, less fluid has a limited densification effect on the samples. Therefore, a higher sintering temperature is required to form a more and widely distributed liquid phase through the reverse reaction of further eutectic reactions in order to obtain near full dense microstructure. At the temperature of 1160 � C, the number of grey phases signifi cantly increased, and the size of grey phases and M6C carbides was similar with each other. Both of the grey phase and the M6C carbide with some pores at the boundary were uniformly distributed in the matrix. As
Fig. 9. EPMA images of the sample sintered at 1160 � C; The values on the top right corner of each figure indicate the top concentration (red). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) 7
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composite powders prepared by high energy ball milling have an average carbide size exceeding 1 μm when sintered at 1160 � C [34]. In addition, the crystal grains do not grow abnormally. By introducing the N element through the powder metallurgy process, a uniform structure can be obtained as well, and the effect of refining the crystal grains can be achieved as in the case of sintering under a nitrogen atmosphere. 3.3. Mechanical properties The hardness values of the sintered samples are given in the line chart of Fig. 11 as a function of temperature. From 1150 to 1200 � C, the hardness tends to be stable as the maximum difference is 2.7 HRC. The hardness values increased significantly from 58.5 HRC at 1150 � C to 62.2 HRC at 1160 � C. Meanwhile the density of samples increased continuously at this stage as shown in Table 3. Hardness is an important performance index of high-speed steel, which can reflect the progress of sintering to some extent [35]. Thus, the hardness of HSS varied with the temperature has a certain consistency with the relative density. As the sintering temperature increased further, the hardness slightly decreased due to the coarsening of carbides as shown in Fig. 7(f). The trend of bend strength with temperature is also shown in the bar chart of Fig. 11. It can be clearly observed that the bend strength of samples rapidly increased from 2502 to 3000 MPa as the sintering temperature increased from 1150 to 1160 � C, because the increasing in sintering temperature accelerated the densification process of the sam ples. However, the bend strength of samples reduced to 2014 MPa when the samples were sintered at 1200 � C. The degradation of bend strength was caused by the coarsening of the microstructure and the formation of secondary pores resulting from the over-sintering of the sample as shown in Fig. 7(e).
Fig. 10. EBSD grain distribution image and inverse pole figure (IPF) orienta tion color scheme of the sample sintered in vacuum atmosphere at 1160 � C. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
4. Conclusions This study revealed that the densification behavior and mechanical properties of M2/VN composites prepared by high-energy ball milling and direct vacuum sintering. On the basis of the performed in vestigations, the following points can be concluded: 1. Refining rate of the composite powder particles decreased gradually with the ball milling time. During the ball milling process, the VN particles were embedded in the matrix phase α-Fe of M2 HSS. After 28 h of ball milling, M6C carbide was continuously refined and ten ded to be lamellar and streamline distribution. 2. Nearly full densification (~99.4% RD) of M2 steel was achieved at 1160 � C by supersolidus liquid phase sintering. At this temperature, the core-shell structure M(C, N) carbonitride with the VN phase as core and MC carbide as shell formed. 3. Addition of VN inhibited the microstructure coarsening and improved both hardness and bend strength. When sintered at 1160 � C, the grain size of the sample is uniform and fine, with an average grain size of only 0.7 μm, and the hardness and bend strength reached 62.2 HRC and 3000 MPa, respectively.
Fig. 11. The hardness (the line chart) and bend strength (the bar chart) of the sintering sample.
Data availability statement
Table 3 Relative density (RD) of specimens sintered at different temperature. T (� C)
1150
1160
1180
1200
1220
RD (%)
95.97
99.35
98.87
98.47
98.18
The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Declaration of competing interest
the MC carbide crystallizes and grows with more stable VN particles as the core. The core-shell grains have the same crystal structure because of the same grain orientation. The interface misfit is small, which cannot be reflected in the EBSD image, indicating that it has a coherent interface with low interface stress. The grain size of the sample is uniform and fine, and the average grain size is only 0.7 μm. Remarkably, the M2/FeV
he authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgements
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Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Dense M2 high speed steel containing core-shell MC carbonitrides using high-energy ball milled M2/VN composite powders Nan Chen, Ren Luo, Huiwen Xiong **, Zhiyou Li * State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, China
A R T I C L E I N F O
A B S T R A C T
Keywords: High speed steel Microstructure High-energy ball milling Sintering Carbonitride
Uniform M2/VN composite powders were prepared via high-energy ball milling of commercial M2 high-speed steel and VN powders, and a promising M2/VN high speed steel with 4 wt% V was obtained by vacuum sin tering. The phases, microstructure of the milled composite powders, and densification behavior, related me chanical properties of sintered steels were investigated. The VN particles were embedded in the matrix phase α-Fe during the ball milling process. After 28 h of ball milling, M6C carbides were continuously refined and showed a streamline distribution. Nearly full densification (~99.4% RD) of M2/VN steel was achieved at 1160 � C by supersolidus liquid phase sintering. Carbonitrides with core-shell structure, exhibiting high N content, were confirmed via the dissolution and precipitation process. The VN particles effectively inhibited the grain growth and improved both the hardness and bend strength. Attracting M2 high speed steel with fine and uniformly dispersed MC carbides was achieved, showing satisfactory bending strength of 3000 MPa and hardness of 62.2 HRC.
1. Introduction 1.1. PM HSS based on the SLPS mechanism High-speed steel (HSS) has been widely used as cutting tools, high load dies, wear-resistant parts due to its excellent mechanical properties [1]. Compared with the conventional cast or wrought parts, HSS fabri cated using powder metallurgy technology (PM HSS) exhibits improved wear resistance, hardness, heat resistance and toughness, owing to the finer and more homogeneous microstructure. In most cases, the densi fication of the sintered PM HSS is achieved through a super-solidus liquid-phase sintering (SLPS) mechanism [2]. The optimum sintering temperature is between the temperature range of the austenite, carbide phase and liquid phase. Low sintering temperature leads to insufficient liquid phase and unsatisfactory density, while high temperature results in the formation of coarser grains and deteriorated performance [3]. General, the main issues of SLPS are as follows: 1) The presence of liquid phase promotes the precipitation of unsatisfactory carbides with mor phologies of rod-like or coarser along the grain boundary [4]; 2) The sintering temperature is relatively high with a narrow window. 3) Careful controlling of the sintering process is needed for achieving the
desired phase components with fine structure [5]. 1.2. M2 HSS with VN particles via vacuum sintering V-containing PM HSS belongs to an attracting material, owing to the in-situ formed fine MC carbide and related high hardness, toughness and wear resistance [6]. As previously reported [7–9], in the nitrogen-rich atmosphere, the sintering temperature of PM HSS with high V content (over 10 wt%) decreases. The N element also results in the enlarged sintering window and the formation of fine carbonitrides and nitrides. The T42 HSS sintered in the nitrogen-rich atmosphere decreased the optimum sintering temperature (OST) from 1220 � C (that of vacuum sintering) to 1160 � C, and widened the sintering window (SW) nearly 30–40 � C [10]. Hence, the preparation of PM HSS with high V content via sintering in nitrogen atmosphere shows a way to handle the above issues of SLPS. Nevertheless, for PM HSS with relatively low V element content, the nitrogen atmosphere sintering is not applicable, owing to the decreased sinterability [7]. The OST of low vanadium high speed steels of M2 HSS (around 2.0 wt% V element content) increased from 1110 � C to 1230 � C when sintered under the nitrogen-rich atmosphere [7]. Based on this, herein, nitrogen was introduced by directly adding
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (H. Xiong), [email protected] (Z. Li). https://doi.org/10.1016/j.msea.2019.138628 Received 12 September 2019; Received in revised form 1 November 2019; Accepted 2 November 2019 Available online 4 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: Nan Chen, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138628
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Fig. 1. Microscopic morphology of the M2 HSS powders (a) and VN powders (b). Table 1 Chemical composition of the M2 HSS powders (mass fraction, %). C
W
Mo
Cr
V
Fe
0.8
6.6
5.2
4.1
1.9
bal.
Table 2 Chemical composition of the VN powders (mass fraction, %). C
V
N
O
3.0
75.6
19.8
bal.
VN particles into the HSS powders, and M2/VN-based HSS with a low V element content of 4 wt% was fabricated via vacuum sintering. The research scheme here may be a promising way to obtain fine MC rein forced PM HSS, and avoid the problem of poor sinterability induced by the nitrogen-rich atmosphere.
Fig. 2. X-ray diffraction patterns of composite powders after milling.
and enhance the mechanical properties of materials [24]. However, there are few reports on the direct sintering preparation process of high-energy ball-milled water atomized HSS powders. The sinterability of M2 HSS is notoriously difficult, and direct sin tering route seems not suitable for it from the past work [25]. Significant improvement of its sinterability is unlikely to be successful through minor modification to its composition or a change to the sintering at mosphere and forming method [26,27]. This paper describes an inves tigation into improving the sinterability of M2 HSS powder by mechanical alloying and to enhance mechanical properties by rein forcement with VN particles. In this research, VN/M2 composite pow ders were prepared using high-energy ball milling. Microstructure evolution of composite powders during the milling process was studied. The densification behavior and mechanical properties of samples are investigated.
1.3. High energy ball milling for PM HSS containing MC Many studies [11–15] have shown that properties of PM HSS were further improved by increasing the amount of carbide forming elements to increase the volume fraction of the hard phases, or by directly adding stable hard phase particles. For instance, the hardness and wear resis tance of M2 HSS are improved by adding carbide particles such as VC [12–14], TiC [16] and NbC [17]. The stability and hardness of MC-type carbide in HSS are higher than those of other carbides such as M6C-type carbide [18], so increasing the content of MC-type carbide is more ad vantageous to improve the hardness and red hardness of materials. When carbide particles are mixed into HSS powder mechanically, the hard phase tends to aggregate between the steel powders, which not only reduces the sintering strength of the materials, but also produces coarse carbides in SLPS [19]. High-energy ball milling has advantages in introducing reinforcing phase, which can refine the powders and disperse the additives evenly into the powder particles [16]. High-energy ball milling generally refers to the process of reducing powder size by mechanical energy [20]. This technology has been suc cessfully applied to synthesize a lot of commercial alloys and composite powders [21–23]. Liu et al. [16]. Refined the powder of M2 HSS by high-energy ball milling and realized the fine dispersion distribution of TiC hard particles in M2 HSS. High energy ball mill can refine the grain structure and make the powders have higher lattice distortion energy. Thus, high-energy ball milling can improve the sinterability of materials
2. Experimental materials and methods 2.1. High-energy ball milling of composite powders The SEM morphology of M2 and VN powders is shown in Fig. 1. The irregular shaped powders with mean particle size of 75 μm were water atomized M2 high speed steel powders, which were provided by €gana €s China Co., Ltd. The VN powders particles are polygonal in Ho shape and they were purchased from Shanghai Yunfu Nanotechnology 2
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Fig. 3. SE SEM images of composite powders milled for (a) 2 h; (b) 4 h; (c) 8 h; (d) 16 h; (e), (f) 28 h.
Co., Ltd. The chemical composition of the M2 HSS powders and VN powders is given in Table 1 and Table 2, respectively. The two kinds of powders were formulated into composite powders with the vanadium element content of 4 wt%. The VN/M2 composite powders were milled by a planetary ball mill with a ball-to-powder weight ratio of 20:1 and a rotating speed of 300 rpm. The ball mill bowls and milling balls are made of stainless steel and GCrl5 alloy, respectively. Argon was used as the protective atmo sphere, and the ball milling time was kept at 2, 4, 8, 16 and 28 h. After the high-energy ball milling, 4.0 wt% of polyethylene glycol, 2.0 wt% of
colloidal graphite and 50 ml of ethanol were added to the composite powders. Polyethylene glycol was dissolved in ethanol before adding into the composite powders. Subsequently, the mixtures were mixed at 150 rpm for 2 h to ensure the adequate homogenization. After ballmilling, drying of the composite powders was carried out in a vacuum oven at 80 � C for 8 h. Finally, the dried composite powders were passed through a 200-mesh sieve for further handling.
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Fig. 4. Internal microstructures of composite powders milled for (a) 2 h; (b) 4 h; (c) 8 h; (d) 16 h; (e), (f) 28 h.
2.2. Preparation of the samples
2.3. Analysis and measurements
The prepared composite powders mixtures were cold-pressed at a pressure of 200 MPa with a dimension of 25 mm � 8 mm � 8 mm. Then, the prepared specimens were put in a tubular hydrogen furnace under flowing H2 gas, and heated at 400 � C for 60 min to remove the binder. The debound samples were placed into the vacuum sintering furnace and first sintered at 1030 � C for 1 h. The purpose of the step above is deoxidize the samples completely. Finally, the sintering of samples was carried out at 1150, 1160, 1180, 1200 and 1220 � C for 1 h, and then cooled to room temperature in the furnace. The heating rate of all the above heating processes was 5 � C/min.
The composite powders and the sintered samples were subjected to phase analysis using a D/Ma2500VB-RA X-ray diffractometer. A FEI Quanta 250 FEG field emission scanning electron microscope (SEM) and a PHOENIX energy dispersive spectrometer (EDS) were used to observe the microstructure and analyze the different phases. Elemental analysis of samples was performed by a JXA-8230 field emission electron probe microanalyzer (EPMA). The grain orientation and size of samples were analyzed using a FEI Helios Nanolab G3 UC electron backscattered diffraction (EBSD) system. A LE COCS-444C/S analyzer and a LECO TC436 N/O analyzer were used to measure carbon and oxygen contents, respectively. Densities of samples were measured by the Archimedes’ 4
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Fig. 5. Internal micro-area N element plane scan of the composite powders. Milled for 28 h; (a) BSED-SEM image of the composite powders; (b) N element dis tribution map.
approximate with each other, their diffraction peaks overlap in Fig. 2. It can be seen from Fig. 3(a) that the shape of VN powders was un changed and the water atomized M2 HSS powder particles were squashed from the original irregular shape after 2 h ball milling. After 4 h of ball milling, the M2 powders became flat and the edges tend to be smooth since the powders were soft at this stage. After 8 h, the powders were obviously broken and refined into fragments, and the degree of edge roundness increased. After 16 h, the powders were further refined, and the number of small particles increased significantly. After 28 h of milling, a large number of fine particles were formed, and the particles tended to be spherical. At the same time, there were aggregates of fine particles. It can be seen from Fig. 3(e) that a large number of fine par ticles were flakily welded on the surface of the large particles, and a depression caused by the peeling of fine particles from the surface of the large particles was observed, indicating that the powders have a mutual balance between crushing and cold welding during ball milling [28]. In order to reveal the phases and morphology of powders, the in ternal microstructures of composite powders are shown in Fig. 4. It can be seen that the initial M2 powder contains white granular M6C carbides and dark grey MC carbides, and most of MC carbides are adjacent to M6C carbides. After 4 h of ball milling, the dark grey block VN was embedded into the M2 matrix, and the MC carbides were separated from the M6C carbides. After 16 h, cracks and pores appeared inside the powder, which may be caused by the crushing and falling off of the hard and brittle phase particles during the milling. In the process of milling for 16 to 28 h, M6C carbides were continuously refined and tended to be lamellar, showing a streamline distribution, and its content was signif icantly reduced. The MC carbides were refined at the initial stage of ball milling, and the change of size and content during 8 h and subsequent ball milling was not obvious, indicating that MC carbides were stable during ball milling [6]. Fig. 5 shows the N-element mapping results of the internal micro-area of the powders after ball milling for 28 h. The N element was dispersedly distributed in M2 HSS, indicating that it was advantageous to prepare a uniform N-containing M2 HSS through the high energy ball milling.
Fig. 6. X-ray diffraction patterns of the sintered samples.
method. Rockwell hardness was provided with a load of 150 kg by a 500MRA-electric hardness tester. Bend strength was measured by three point bending with a span of 14.5 mm. Each numerical value was repeated at least three times and the results were expressed using the mean value. 3. Results and discussion 3.1. Phases, and microstructure of the M2/VN powders after milling Fig. 2 shows the X-ray diffraction patterns of the composite powders with different milling times. The initial powder mixtures mainly include α-Fe, M6C and M (C, N). As the ball milling time increased, the diffraction peaks were significantly broadened and weakened due to the refined crystal lattice and internal stress. In the middle and late stages of ball milling, some of the M6C and M (C, N) peaks disappeared. It is noteworthy that the change of diffraction peaks of carbides is not obvious because of their small volume fraction. The enlarged picture of the diffraction peaks between 42� and 45� is shown in Fig. 2. The first strong peak of Fe gradually shifted to the left, indicating that W, Mo and other large atoms of M6C were continuously dissolved into α-Fe. Since VC and VN are both cubic structures and the cell parameters are
3.2. Microstructure evolution during sintering XRD patterns of sintered samples are presented in Fig. 6. The M2/VN alloyed samples show obvious diffraction peaks of M6C and M (C, N) after sintering at 1150 � C. The reason is that the lattice distortion is eliminated after sintering and the diffraction peak becomes sharper [29]. Compared with the sample sintered at 1150 � C, the diffraction peak of Fe is obviously weakened and broadened after sintering at 5
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Fig. 7. BSED-SEM images of samples sintered in vacuum atmosphere at (a), (b) 1150 � C; (c), (d) 1160 � C; (e) 1180 � C; (f) 1200 � C.
1160 � C and above. While the peaks of M6C become weaker, the peaks of MC become stronger, and some peaks of M6C disappear, indicating that the content of M6C decreases and the content of MC increases. Mean while, the diffraction peak of α-Fe shifted to the higher angles. This is because the interplanar distance decreased due to dissolution of large atomic radius alloy elements in the matrix [30]. The microstructure evolution of the samples sintered at different temperatures is shown in Fig. 7. When the samples were sintered at 1150 � C, Closed circular pores remained at the boundary between the carbides and matrix as shown in Fig. 7(a) and (b). The size of M6C carbides (white phase) was significantly larger than that of M (C, N) and MC (grey phase). In most cases, High speed steel densifies via
supersolidus liquid phase sintering (SLPS). For supersolidus liquid phase sintering, sufficient liquid is required to enhance the densification of samples [16]. However, due to the gravity of samples, excess liquid can cause loss of slump and dimensional control. For supersolidus liquid phase sintering of high-speed steel, usually the amount of liquid as well as the densification rate and microstructure changes is controlled by the sintering temperature typically. The liquid phase has greater chance of appearing firstly at the interfaces between the enhanced particles and the matrix [16]. Carbide reinforcements and γ-Fe matrix form liquid phase through the reverse of eutectic reaction at the interfaces (γ-Fe þ MxCy→L) [31]. Some liquid phase usually migrates along the surface of γ-Fe matrix particles to the sintered neck, while the γ-Fe matrix that does 6
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shown in Fig. 7(d), the grey phases corresponded to a vanadium-rich carbonitride formed as core-shell structures with the VN phases as core and MC carbides as shell [32]. The shell formed through the dissolution and precipitation of original carbides in the M2 HSS [31]. When the temperature rose to 1180 � C, the carbides grew up signifi cantly. Many closed circular pores formed at the interfaces between carbides and the matrix. In addition, the carbides with core-shell structure have largely disappeared. At the temperature of 1200 � C, abnormal growth of MC carbide was observed. More liquid phase accelerated the rate of atomic diffusion. Through the dissolution-precipitation mechanism, smaller carbide dissolved in the liquid phase and then precipitated on the surface of larger carbide. The appearance of crescent carbides was the evidence of massive liquid phase in sintering process [33]. The graph of the average M6C carbide size as a function of temperature is shown in Fig. 8. The size of M6C carbide continued to increase from 0.89 μm to 1.38 μm along with the rising of the temperature. Fig. 9 shows the EPMA images of the sample sintered at 1160 � C. M (C, N) phase has high vanadium content in the central region and low vanadium content in the surrounding area. In the central region of M (C, N) phase, there are nitrogen-rich regions which are in red color. In the grey M (C, N) carbide, there is a gradient distribution of elements, which is corresponding to the reported non-uniform distribution of the MC phase in fully alloyed HSS prepared by gas atomization-hot isostatic pressing [30]. This also indicates that carbonitrides having a core of VN formed during the sintering process. Fig. 10 shows the EBSD grain distribution image and inverse pole figure (IPF) orientation color scheme of the sample sintered at 1160 � C. Black regions are caused by long-term vibration polishing resulting in carbide bulge. The M6C carbide is mostly clumpy and the MC carbide is mostly near-spherical as shown in Figs. 10 and 7(d). This reason may be that the M6C carbide directly precipitates from the liquid phase, while
Fig. 8. Average M6C carbide size of the samples at different temperatures.
not react with carbide reinforcements is still solid phase sintering. Thus, less fluid has a limited densification effect on the samples. Therefore, a higher sintering temperature is required to form a more and widely distributed liquid phase through the reverse reaction of further eutectic reactions in order to obtain near full dense microstructure. At the temperature of 1160 � C, the number of grey phases signifi cantly increased, and the size of grey phases and M6C carbides was similar with each other. Both of the grey phase and the M6C carbide with some pores at the boundary were uniformly distributed in the matrix. As
Fig. 9. EPMA images of the sample sintered at 1160 � C; The values on the top right corner of each figure indicate the top concentration (red). (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.) 7
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composite powders prepared by high energy ball milling have an average carbide size exceeding 1 μm when sintered at 1160 � C [34]. In addition, the crystal grains do not grow abnormally. By introducing the N element through the powder metallurgy process, a uniform structure can be obtained as well, and the effect of refining the crystal grains can be achieved as in the case of sintering under a nitrogen atmosphere. 3.3. Mechanical properties The hardness values of the sintered samples are given in the line chart of Fig. 11 as a function of temperature. From 1150 to 1200 � C, the hardness tends to be stable as the maximum difference is 2.7 HRC. The hardness values increased significantly from 58.5 HRC at 1150 � C to 62.2 HRC at 1160 � C. Meanwhile the density of samples increased continuously at this stage as shown in Table 3. Hardness is an important performance index of high-speed steel, which can reflect the progress of sintering to some extent [35]. Thus, the hardness of HSS varied with the temperature has a certain consistency with the relative density. As the sintering temperature increased further, the hardness slightly decreased due to the coarsening of carbides as shown in Fig. 7(f). The trend of bend strength with temperature is also shown in the bar chart of Fig. 11. It can be clearly observed that the bend strength of samples rapidly increased from 2502 to 3000 MPa as the sintering temperature increased from 1150 to 1160 � C, because the increasing in sintering temperature accelerated the densification process of the sam ples. However, the bend strength of samples reduced to 2014 MPa when the samples were sintered at 1200 � C. The degradation of bend strength was caused by the coarsening of the microstructure and the formation of secondary pores resulting from the over-sintering of the sample as shown in Fig. 7(e).
Fig. 10. EBSD grain distribution image and inverse pole figure (IPF) orienta tion color scheme of the sample sintered in vacuum atmosphere at 1160 � C. (For interpretation of the references to color in this figure legend, the reader is referred to the Web version of this article.)
4. Conclusions This study revealed that the densification behavior and mechanical properties of M2/VN composites prepared by high-energy ball milling and direct vacuum sintering. On the basis of the performed in vestigations, the following points can be concluded: 1. Refining rate of the composite powder particles decreased gradually with the ball milling time. During the ball milling process, the VN particles were embedded in the matrix phase α-Fe of M2 HSS. After 28 h of ball milling, M6C carbide was continuously refined and ten ded to be lamellar and streamline distribution. 2. Nearly full densification (~99.4% RD) of M2 steel was achieved at 1160 � C by supersolidus liquid phase sintering. At this temperature, the core-shell structure M(C, N) carbonitride with the VN phase as core and MC carbide as shell formed. 3. Addition of VN inhibited the microstructure coarsening and improved both hardness and bend strength. When sintered at 1160 � C, the grain size of the sample is uniform and fine, with an average grain size of only 0.7 μm, and the hardness and bend strength reached 62.2 HRC and 3000 MPa, respectively.
Fig. 11. The hardness (the line chart) and bend strength (the bar chart) of the sintering sample.
Data availability statement
Table 3 Relative density (RD) of specimens sintered at different temperature. T (� C)
1150
1160
1180
1200
1220
RD (%)
95.97
99.35
98.87
98.47
98.18
The raw/processed data required to reproduce these findings cannot be shared at this time as the data also forms part of an ongoing study. Declaration of competing interest
the MC carbide crystallizes and grows with more stable VN particles as the core. The core-shell grains have the same crystal structure because of the same grain orientation. The interface misfit is small, which cannot be reflected in the EBSD image, indicating that it has a coherent interface with low interface stress. The grain size of the sample is uniform and fine, and the average grain size is only 0.7 μm. Remarkably, the M2/FeV
he authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
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Acknowledgements
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