Full Terms & Conditions of access and use can be found at https://www.tandfonline.com/action/journalInformation?journalCode=ypom20 Powder Metallurgy ISSN: 0032-5899 (Print) 1743-2901 (Online) Journal homepage: https://www.tandfonline.com/loi/ypom20 In-situ VN reinforced powder metallurgy M30 steels prepared from water atomized powers via pressureless sintering Haixia Sun, Fang Yang, Qian Qin, Biao Zhang, Alex A. Volinsky & Zhimeng Guo To cite this article: Haixia Sun, Fang Yang, Qian Qin, Biao Zhang, Alex A. Volinsky & Zhimeng Guo (2020) In-situ VN reinforced powder metallurgy M30 steels prepared from water atomized powers via pressureless sintering, Powder Metallurgy, 63:1, 43-53, DOI: 10.1080/00325899.2020.1717075 To link to this article: https://doi.org/10.1080/00325899.2020.1717075 Published online: 26 Jan 2020. Submit your article to this journal Article views: 8 View related articles View Crossmark data
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Full Terms & Conditions of access and use can be found athttps://www.tandfonline.com/action/journalInformation?journalCode=ypom20
In-situ VN reinforced powder metallurgy M30 steels prepared from wateratomized powers via pressureless sinteringHaixia Suna, Fang Yanga, Qian Qina, Biao Zhanga, Alex A. Volinsky b and Zhimeng Guoa
aInstitute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing, People’s Republic of China;bDepartment of Mechanical Engineering, University of South Florida, Tampa, FL, USA
ABSTRACTIn this paper, in-situ VN reinforced powder metallurgy high-speed steel was fabricated fromwater atomised powder via pressureless sintering. During sintering, N would replace C toreact with V, resulting in the formation of in-situ VN phase. According to the first-principlescalculations, the formation energy of VN was – 9.45 eV while that of VC was – 9.08 eV. Therelative density of as-sintered samples reached up to 99%. Homogeneous microstructure andfine reinforced phases were obtained. Ultrafine VN phase (0.5 μm) was uniformly distributedin the matrix. As a result, the mechanical properties were improved. The optimal mechanicalperformance was obtained in the sample with 1.0 wt.% C addition. The hardness, bendstrength and impact energy were 65 HRC, 3011 MPa and 18–22 J, respectively. Besides, thesintering window significantly increased from ∼10 to 30∼50°C.
ARTICLE HISTORYReceived 7 October 2019Revised 5 January 2020Accepted 9 January 2020
Owing to their excellent strength, toughness and hard-ness, tool steels are widely used in cutting tools andwear parts. Among tool steels, high-speed steel (HSS)is one of the most common kinds with special micro-structure and excellent properties. In terms of themicrostructure, fine carbides are uniformly distributedin the steel matrix. Thereinto, the carbides character-istics, such as particle size, distribution, structure andinterface with the matrix play an important role indetermining the microstructure and mechanical prop-erties of the HSS.
It is known that the carbides’ morphology of HSSdepends on the fabrication process [1–3]. Besides, thecarbides structure and interface character are closelyrelated to the phase composition. Carbide, particularlyprimary carbide, is crucial, which significantly affectsthe strength, hardness, and processability. The primarycarbide mainly consists of tungsten and molybdenum-rich M6C, vanadium-rich MC, and even M2C with agranular or irregular shape in the cast HSS [4–6]. Gen-erally, the preparation process mainly includes castingand powder metallurgy (PM). The M2C is a representa-tive carbide in cast steels with relatively large and bulkierparticle size. In order to refine the coarse carbidesformed during the casting process, a large amount offorging deformation is required, resulting in the strip-shaped microstructure which causes the anisotropyproperties [7]. The PM process can solve the problemsof coarse primary carbide and material anisotropy.
In general, the PM process for preparing HSS mainlyincludes hot isostatic pressing (HIP), pressureless sinter-ing, and injection moulding. During the HIP process, gasatomisedHSS powder and steel capsule are employed [8].Besides, the requirements for HSS powder quality arevery stringent and themanufacturing process is relativelycomplex. These greatly restrict the development of PMHSS. It is well known that gas atomised powder is spheri-cal with extremely low oxygen content, while water ato-mised powder is mainly irregular with relatively highoxygen content. During the gas atomisation process, theheat exchange of the particles in the gas environment isslow, resulting in a low particle cooling rate. In suchcase, the particle droplets automatically shrink into thesphere. In the contrary, because of the chilling effect ofwater on the atomised droplets, the particles solidifyinto alloy powder instantaneously, resulting in the lesssurface tension of water atomised powder with an irregu-lar shape. Therefore, gas atomised HSS powder is moresuitable for HIP sintering with low oxygen content. Asto pressureless sintering, it is a promising method to pro-duce HSS using water atomised powder. Generally, wateratomised HSS powder is mainly adopted, which can besintered to nearly full densification. Besides, the sinteringis carried out under vacuum or nitrogen atmosphere. Atpresent, the densification is achieved by supersolidusliquid phase sintering (SLPS) [9,10]. In such case, densifi-cation and microstructural control are correlated withpseudobinary phase diagrams. The optimal sinteringtakes place on heating to a four phase austenite +M6C
CONTACT Fang Yang [email protected] Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing100083, People’s Republic of China; Zhimeng Guo [email protected] Institute for Advanced Materials and Technology, University of Science andTechnology Beijing, Beijing 100083, People’s Republic of China
+MC+ liquid region.When the solidus-liquidus intervalis wider and the slope of the solidus line is deeper, the sin-tereability is superior. Therefore, many efforts have beenreported focusing on the preparation of HSS with wateratomised powders. The studies on how to achieve pres-sureless sintering of HSS have received extensive atten-tion due to their low cost [11,12]. Giménez et.al studiedthe microstructure characteristics of HSS fabricated byHIP, vacuum sintering, and nitrogen sintering [13].After sintering under different nitrogen partial pressureconditions, there was no big difference in the M6Cphase, except for the MC phase. The transformation ofstiffer MN phase from the MC phase occurred, whichhad a beneficial effect on toughness and strength ofHSS. As a drawback, the reaction was accompanied bythe formation of the M3C network carbides, which wasdetrimental for achieving good performance. A similarreport was presented in Ref. [14]. Although fine andspherical MN phase was obtained in the V-containingHSS, theM3C network carbides and coarseM6C carbideswere also formed, resulting in the destruction of isotropicproperties [14]. Although TiC, NbC, andMnS reinforcedPM HSS have been studied [15–17], there is no relatedreports on preparing MN-reinforced HSS with isotropicproperties.
In this paper, in-situ VN reinforced PM HSS wasprepared by pressureless sintering under the nitrogenatmosphere. Among commonly used pressureless sin-tering processes, nitrogen sintering was widelyemployed in the preparation of metal and ceramicparts [18–20]. As opposed to the common pressurelesssintering process via adding graphite to balance thecarbon content, this study was designed to decreasethe carbon content in the V-containing HSS. Despitethat, the PM HSS with VN strengthening possessedexcellent performance. The carbon effects and sinteringbehaviour in the PM HSS were studied. Besides, themicrostructure, elemental distribution and phase com-position of the VN reinforced phase were investigated.Accordingly, the hardness, bend strength and impactenergy were measured. Based on the analysis of phaseformation energy and standard Gibbs free energy, theformation mechanism of carbides and nitrides in theHSS was also clarified.
2. Experimental procedures
2.1. Processing
The raw material was water atomisation HSS powderwith a nominal composition of 4.33Cr–6.30W–4.96Mo–2.97V–8.72Co–xC-bal. Fe (x = 0.8, 1, and 1.2wt.%). The powder was purchased from the Antai Tech-nology Co., Ltd., with an average powder particle size of10 μm. The preparation process comprised the followingsteps. First, the powder was compressed by the hydraulicpress with a pressure of 700 MPa. Then, the compacts
were pressureless sintered at different temperaturesunder nitrogen atmosphere with a heating rate of 5°C min–1. The sintering temperature was chosen accord-ing to the study reported in Ref. [13], and the gas flowrate was 300 mL min–1. Subsequently, furnace coolingto room temperature was performed. Lastly, heat treat-ment was carried out as follows: austenitising at 1150°Cfor 10–20 min, oil quenching, and triple tempering at500–600°C for 1–2 h. For comparison, traditional M30samples were also prepared by vacuum sintering under10−1–10−3 Pa. The heat treatment parameters were thesame as before.
2.2. Characterisation
The density of the as-sintered samples was measuredby the Archimedes method. Five samples for each pro-cessing condition were employed to confirm repeat-ability. Mechanical properties were evaluated in termsof the hardness, bend strength, and room temperatureimpact tests. The hardness of the samples after heattreatment was determined by a TH320 Rockwell hard-ness tester. The bend strength was measured by thethree-point bending test with an electronic universaltesting machine (CMT6140) at a loading rate of 0.5mm s–1. The size of the tested samples was 35 mm ×5 mm × 5 mm with a span of 30 mm. Besides, theimpact energy was measured by the non-notch impacttest using a CMT4105 electronic universal testingmachine, in which the sample size was 55 mm × 10mm × 10 mm. The abrasive wear resistance test wascarried out under load of 30 N for 120 s, using a ML-100 wear testing machine. Abrasive was 600# SiC sand-paper and the sample size was φ10 mm*15 mm. Theoxygen and nitrogen contents were measured using apulsed infrared thermal conductivity oxygen-nitrogenanalyser (ONH-2000). The phase formation energywas calculated by the norm-conserving pseudo-poten-tial method (CASTEP). Powder morphology, micro-structure, fracture and worn surface were analysedusing a scanning electron microscopy (SEM, PhilipsLEO-1450) equipped with energy-dispersive spec-troscopy (EDS, Model JEOL, JEOL Ltd., Tokyo,Japan). Element distribution was observed by a fieldemission scanning electron microscope (FESEM,Zeiss Supra55). Phase identification was analysedusing a transmission electron microscope (TEM, Tec-nai G2 F30 S-TWIN) operating at 300 kV. Phase analy-sis was carried out using an X-ray diffraction (XRD,Shimadzu XRD-6000, Cu Kα target, 40 kV and 40 mA).
3. Results and discussion
3.1. Densification
The morphology of raw HSS powders is shown inFigure 1. The powders were irregular and ellipsoid.
44 H. SUN ET AL.
Fine powder (∼10 μm) was employed to ensure the sin-terability in the pressureless sintering process. Owingto the subtle distinctions in powder composition, thefour kinds of green compacts were sintered at differenttemperatures to guarantee the as-sintered samplesobtaining the optimal mechanical performance.Accordingly, the temperature parameters and sinteringbehaviour of different samples are summarised inTable 1. As reported in Ref. [13], the optimal sinteringtemperature (OST) is the lowest sintering temperaturefor achieving the relative density of above 98%. Thesintering window (SW) is corresponding to the temp-erature interval between the OST and the incipientmelting temperature of the M6C phase. The higherthe SW value, the lower the probability of grain growthcaused by local overheating. As presented in Table 1,with the carbon content decreasing, the density of thesamples sintered under nitrogen decreased, while theOST increased. The relative densities of all sampleswere over 98%. Besides, the relative density of theM30 sample was about 99%, which was approximatelyequal to the value of the NM30-2 sample. However, itwas worth noting that the SW values of the NM30-2and NM30-3 samples were obviously higher than theM30 sample.
3.2. Microstructure and properties
The microstructures of different as-sintered samplesare shown in Figure 2. A typical HSS microstructurewith reinforced carbide or nitride phases distributedin the matrix was observed. The dark grey phase wascorresponding to the matrix phase. In the M30 sample,two kinds of reinforced carbide phases were obtained,
as shown in Figure 2(d). According to the EDS results,the white phase was corresponding to the M6C carbide,enriched with tungsten and molybdenum elements.The light grey phase was the V-rich MC carbide.These results were consistent with the literatures[21,22]. Different from the M30 sample, there wasanother black phase formed in the samples sinteredunder nitrogen, as shown in Figure 2(a–c). It was ahard phase enriched with V and N elements. Asshown in Figure 2, the average particle size of carbideswas about 1–2 μm, except for the NM30-1 sample.Lath-shaped carbides were also found in Figure 2(a).It was worth noting that the formed nitride phaseswere ultrafine, less than 1 μm. The microstructure ofsamples sintered under nitrogen was distinctly differ-ent from the common PM HSS [23].
To further confirm the phase composition and crys-tal structure, FESEM, XRD and TEM investigationswere performed. The element distribution of Fe, V,Mo, W, Co, C, and N elements is presented in Figures3 and 4. Fe and Co were thought to be dominant in thematrix while C was distributed uniformly in the matrixand the reinforced phase. It can be observed that N andV were mainly precipitated in the black region, corre-sponding to the nitride phases. As for the N-free brightphase, it mainly consisted of W and Mo elements. Thedistribution of W and Mo elements were completelyoverlapped. Therefore, both the M6C and nitridephases were observed in all NM30 samples, and noMC phase was detected. As reported in Ref. [13], thechange in the sintering process would affect the for-mation of the MC phase when the HSS samples weresintered under nitrogen. The formation mechanismwould be discussed in further detail.
Figure 1. SEM images of raw HSS powder at: (a) low, and (b) high magnification.
Table 1. C, N and O contents and density for HSS samples sintered at optimum sintering temperatures (OST).
Figure 2. Microstructures of different as-sintered HSS samples: (a) NM30-1, (b) NM30-2, (c) NM30-3, and (d) M30.
Figure 3. Fe, V, Mo, W, Co, C and N elements distribution in the NM30-1 HSS sample.
Figure 4. Fe, V, Mo, W, Co, C and N elements distribution in the NM30-2 HSS sample.
46 H. SUN ET AL.
The phase composition was characterised by XRD,as shown in Figure 5. Overall speaking, M6C and MCwere detected in the M30 sample, while ferrite, M6C,and MN were detected in the NM30 sample. It wasworth noting that M2C was also observed in theNM30-1 sample, which would be harmful for themechanical performance. The only difference betweenthe NM30-1 and NM30-2 samples was the carbon con-tent. Compared with the NM30-2 sample, an extra 0.2wt.% C was introduced into the NM30-1 sample. It canbe inferred that the carbon content should be strictlycontrolled. To further analyse the phases character-istics, the peak details were detected from 31° to 44°and 43° to 46°, as shown in Figure 5(b,c). Face-centredcubic MC phase was only observed in the M30 sample.The MN phase was only found in the samples sinteredin nitrogen. No impurity phases, like M2C and V2N,were detected in the NM30-2 sample compared withthe NM30-1 sample. For the MC and MN phases,nitrogen and carbon atoms were occupied in the octa-hedral spaces of the vanadium atom. Thereinto, theradius of the nitrogen atom (0.071 nm) was smallerthan the carbon atom (0.77 nm). As to the ferrite,there was no peak position shift between the M30and NM30-2 samples. However, for the NM30-1sample, the peak of ferrite shifted to a higher diffractionangle. This implied that there were atoms dissolved inthe matrix, leading to the decreasing of lattice spacing.It can be inferred that when sintered under nitrogen, Nwould react with V to form the VN phase. Comparingthe M30 and NM30-1 samples, although the carboncontent was the same, V atoms were occupied by Natoms, leading to an excess carbon content. Accordingto the XRD analysis, part of the excess C would dissolveinto the matrix, resulting in the peak shift, and anotherpart would form the M2C phase. Therefore, thelath-shaped carbides, shown in Figure 2(a), might bethe M2C phase enriched with molybdenum and
tungsten, which were similar to the carbides in castHSS [24,25].
The reinforced phase was further identified by TEManalysis, as presented in Figure 6. From EDS resultsshown in Figure 6(a,b), region ‘A’ was the carbidephase enriched with W and Mo, while region ‘B’ wasthe nitride phase enriched with V. The correspondingSAED patterns are shown in Figure 6(d,e). TheSAED pattern of region ‘A’ was indexed as cubicM6C phase with a = 1.114 nm. The zone axis was[001]. Another SAED pattern revealed that region ‘B’was corresponding to face-centred cubic VN phasewith a = 0.414 nm. The zone axis was also identifiedas [001]. Besides, the grain size of the M6C phase wasabout 1 μm. However, the grain size of the VN phasewas much smaller, about 0.5 μm. Fine VN reinforcedphase was formed during sintering, which mayenhance the mechanical performance.
Accordingly, the hardness, bend strength andimpact energy of different samples are listed inTable 2. The hardness, bend strength and impactenergy of the M30 sample were 65 HRC, 2962 MPaand 15–20 J, respectively. The optimal mechanical per-formance was obtained in the NM30-2 sample. Therewas no big difference in hardness. However, the bendstrength of the NM30-2 sample was up to 3011 MPa,which was 50% higher than that of the NM30-1sample. The impact energy increased from 2–5 to18–22 J. Figure 7 presents the mechanical propertiesof different PM HSS samples. Comparing the M30and NM30-2 samples, there was an obvious yield pla-teau observed in NM30-2 sample, just like the greenmark in Figure 7. As a result, the strain value ofNM30-2 was much larger than M30, indicating thatfine VN phase could contribute to improving the plas-ticity of HSS parts. It can be inferred that the failure oftool tipping would be significantly reduced, resulting ina prolonged service life of HSS. The mechanical
Figure 5. XRD patterns of the as-sintered HSS samples detected at: (a) 20–90°, (b) 31–44°, and (c) 43–46°.
POWDER METALLURGY 47
properties were strongly related to the microstructure.Therefore, the performance enhancement was mainlyattributed to the formation of fine VN reinforcedphase.
The fracture morphology of different HSS samples isshown in Figure 8. The fracture mainly occurred alongthe bonding interface between the matrix and thereinforced phases. It can be clearly observed thatthere were larger reinforced phases in the NM30-1sample, indicated by the white arrow in Figure 8(d).These might be the lath-shaped carbides. The largeblock M6C or M2C carbides caused considerably highstress concentration in the surrounding area, leading
to early bend or fracture. Hence, the bend strengthand impact energy of the NM30-1 sample were muchlower. In contrast, the NM30-2 sample had homo-geneous microstructure with fine carbides and nitridesevenly distributed in the matrix, which was beneficialfor withstanding much larger bend and impact forces.Therefore, the mechanical performance of the NM30-2sample was much higher than NM30-1.
Moreover, the wear resistance analysis was also car-ried out, as shown in Figure 9. It can be seen thatNM30-1 sample showed the worst wear resistance.Obvious plough scratch was observed in Figure 9(a).The grooves were relatively deeper. In contrast, thewear resistance of NM30-3 sample was much betterthan that of NM30-1 sample. It might be attributedto fine hard particles existed in the matrix. Overall,NM30-2 and M30 samples had the best wear resist-ance. The worn surface for NM30-2 sample was similarto M30 sample, as shown in Figure 9(f,h). As shown inFigure 2, the carbides and nitrides in NM30-2 and M30samples were fine and homogeneous, and the particle
Figure 6. EDS analysis of region (a) A, (b) B, (c) TEM image and SAED patterns of region (d) A, and (e) B in sample NM30-2.
Figure 7. (a) Force-displacement curve, and (b) stress–strain curve of PM HSS samples.
Table 2. Mechanical properties of PM HSS.Carbon content
spacing was relatively small, resulting in the obtain-ment of better wear resistance. Besides, according tothe wear test shown in Figure 9, VC and VN shouldhave similar wear resistance.
3.3. Formation mechanism
To understand the formation mechanism, the standardreaction Gibbs free energy of VN and VC was calcu-lated, as shown in Figure 10. The reaction can be pre-sented as Equation (1). At 1210°C, the ΔG values of VCand VN were −85.736 and −471.58 J, respectively. TheΔG of VN was much more negative than VC, indicat-ing that the formation of the VN phase had a higher
tendency. As mentioned in Table 1, the sintering temp-erature of the HSS samples was about 1200°C. There-fore, nitrogen gas would react with V to form in-situVN phase during sintering.
V+ C(N) = VC(N) (1)
Furthermore, the first principles calculations wereemployed to calculate the phase formation energy.The crystal models of VC/VN were built, as shown inFigure 11(a). The VC/VN phase had a simple face-centred cubic structure. Accordingly, the formationenergy of VC and VN phase was calculated, as shownin Figure 11(b). The lower the formation energy, the
Figure 8. Fracture morphology of different HSS samples: (a, b) NM30-2, (c, d) NM30-1.
Figure 9. SEM images of worn tools: (a, e) NM30-1, (b, f) NM30-2, (c, g) NM30-3, and (d, h) M30.
POWDER METALLURGY 49
more stable the phase was, and correspondingly, thehigher its formation tendency. It can be observed thatwith the N content increase in the VC phase, the for-mation energy value was more negative. The valuesof the VC and VN phase were −9.08125 and−9.44895 eV, respectively. This further proved thestability of the VN phase was higher than VC, whichwas consistent with the result shown in Figure 10.Compared with VC, VN had more negative ΔG andphase formation energy. During the nitrogen sinteringprocess, in-situ VN reinforced phase would form andits particle size was ultrafine.
The driving force of sintering was the reduction offree energy (ΔZ) of the sintering system. As shown inTable 1, the density of the NM30-2 sample was 7.98g cm–3, which was slightly higher than the M30 sample.It can be inferred that the ΔZ value would go downmore when forming VN rather than VC, leading tothe density further being closer to the theoretical den-sity. Besides, nitrogen gas provided an activated sinter-ing atmosphere for V-containing HSS. Generalspeaking, in the final stage of sintering, the eliminationof pores depended on the difference between residualgas pressure and surface tensile stress. As presented
in Equation (2), the Ps represented the force to promotedensification sintering. Thereinto, the Pv was theresidual gas pressure in the pores, while
2gr
was corre-sponding to the surface tensile stress. If the
2gr
valuewas larger than Pv, the pores can keep shrinking. Incontrast, if the gas remained in the sample, thePv value would be large enough to exceed the
2gr,
resulting in the stop of sintering shrinkage and theexistence of pores [26]. When sintered under nitrogen,the residual gas was nitrogen, which would diffuse intothe powders to form nitride. As a result, the sinteringdriving force significantly improved. Therefore, thenitrogen activated sintering was a promising methodto prepare full density V-containing HSS parts.
Ps = 2gr− Pv (2)
3.4. Discussion
The above results indicated that the in-situ formednitride opened new possibilities for fabricating HSSwith high performance. The sintering densificationwas decided by the sintering neck forming and poreeliminating. Fine HSS powder (10 μm) had a relativelylarge surface area and energy, resulting in the obtain-ment of high relative density (>98%) when pressurelesssintered under vacuum or nitrogen atmosphere. Whensintered under nitrogen, the formation energy of VNwas more negative, resulting in the formation of in-situ fine VN phase instead of VC phase. Besides, nitro-gen effect on improving the sinterability of HSS can bedescribed in terms of the enlargement of the austenite+M6C +MC + liquid region. In the V-containing HSS,the relationship between C content and N absorptionplayed an important role in controlling densificationsintering. When sintering under nitrogen, VN wasformed instead of VC. In such case, free carbon wasavailable to dissolve into the matrix, thus decreasingthe solidus temperature. Furthermore, as the SWvalue of the NM30-2 sample greatly increased, the
Figure 10. Standard Gibbs free energy change (ΔG°) as a func-tion of temperature.
Figure 11. (a) Crystal structure of VC1−xNx (x = 0–1), and (b) corresponding formation energy change as a function of nitrogenatomic percent.
50 H. SUN ET AL.
difficulty in controlling the sintering process signifi-cantly reduced, which would make industrialised pro-duction more convenient and reduce therequirements for production equipment.
The composition of the reinforced phase was closelyrelated to the carbon content. As the carbon contentdecreased, the M2C phase gradually disappeared andthe M6C amount slightly reduced. When the carboncontent was 1.0 wt.%, the reinforced phases were justM6C and MN. Homogenous microstructure wasobtained with fine M6C (1 μm) and ultrafine MNphase (0.5 μm) distributed in the matrix. Accordingto TEM analysis, M6C phase had a complex cubicstructure. There were 96 metal atoms (mainly Fe, W,and Mo) and 16 carbon atoms in the unit cell. Rather,MN phase possessed a simple face-centred cubic struc-ture, in which M atoms occupied the top and centresites, while N atoms occupied the octahedral gaps.Because the radius of nitrogen atom (0.071 nm) wasless than the carbon atom (0.077 nm), the value ofRn/Rm was smaller than Rc/Rm. Therefore, the formedMN was more stable than the MC [27].
As shown in Table 1, there was an obvious differ-ence on the C and N contents in the as-sintered HSSsamples. The transformation of the reinforced phasein HSS is sketched in Figure 12. Figure 12(a) presentsa carbon balance state in HSS when sintered under vac-uum. When sintered under nitrogen, if the HSS systemwas in the carbon imbalance state, lath-shaped M2Cand M6C were formed with larger particle size. Forthe M30 sample, the balanced carbon content was 1.2wt.%. Although the carbon content was the same inthe NM30-1 sample, the formation of the VN phasewould lead to an excess of carbon. This was attributedto the lower formation energy of VN. Therefore, thecarbon content should be properly decreased, justlike the NM30-2 sample. A carbon–nitrogen balancestate was essential for obtaining high-performanceHSS.
This new HSS was designed with lower carbon con-tent to achieve the balance between carbon and nitro-gen. During sintering, nitrogen diffused into thematrix to take part in the reaction of nitride formation.The nitride was predominantly the VN phase, whichhad lower standard reaction Gibbs free energy and
more negative phase formation energy. Comparedwith VC, in-situ VN was smaller and more stable, lead-ing to higher strength and toughness. Therefore, theoptimal mechanical performance was obtained in theNM30-2 sample. The hardness, bend strength andimpact energy were 65 HRC, 3011 MPa and 18∼22 J,respectively.
It is well known that these materials do not fracturein a classical Griffith manner, but have an intermediatecrack growth stage in monotonic fracture, and alsoshow R-curve behaviour [28]. Typical plastic zonesizes can be calculated from the relationship: radiusof the plastic zone Rp = 1/6p[KIC/syld]
2, where KIC
is the fracture toughness, syld is the yield stress. In gen-eral, syld ≈ (3.27Hm)MPa, and Hm is the micro-hard-ness of the steel matrix [15]. In simple terms, Rp
would increase with the increase of KIC. From thebend strength curve shown in Figure 7, the intermedi-ate crack growth stage of NM30-2 sample was moreslow, indicating that it might potentially have a highfracture toughness. According to the radius of plasticzone, it could be inferred that NM30-2 sample mayhave larger plastic zone. However, this need to beinvestigated further and in depth. In our futureresearch, the fracture behaviour and cutting tool per-formance of VN reinforced HSS would be studied indetail.
4. Conclusions
In this study, in-situ VN reinforced PM HSS was pre-pared from water atomised powders via pressurelesssintering. The main conclusions are as follows:
(1) Fine raw powder was employed to prepare HSSwith above 98% of theoretical density. When sin-tered under nitrogen, a carbon–nitrogen balancestate was essential to obtain the M6C/VNreinforced HSS.
(2) Better comprehensive mechanical properties wereobtained in the NM30-2 sample containing 1.0wt.% C. The hardness, bend strength and impactenergy were 65 HRC, 3011 MPa and 18–22 J,respectively.
Figure 12. Schematics of microstructure development in HSS: (a) carbon balance state, (b) carbon imbalance state, (c) new carbon–nitrogen balance state.
POWDER METALLURGY 51
(3) Homogenous microstructure was obtained withfine M6C and ultrafine VN uniformly distributedin the matrix. The average grain size of VN was0.5 μm, which was much smaller than VC, resultingin the enhancement of mechanical performance.
(4) The formation mechanism was elucidated. Thestandard Gibbs free energy of VN was muchlower than the VC phase. Besides, the formationenergy of VN was −9.45 eV, while that of VCwas −9.08 eV.
Disclosure statement
No potential conflict of interest was reported by the authors.
Funding
This work was supported by the China PostdoctoralResearch Foundation [No. 2018M641188] and the Funda-mental Research Funds for the Central Universities [No.FRF-TP-18-025A1].
Notes on contributors
Haixia Sun is Ph.D. candidate in Engineering at the Univer-sity of Science and Technology Beijing. She is engaged in theresearches on powder metallurgy iron-based materials.
Fang Yang is a Ph.D. in Engineering and works as a lecturerat the University of Science and Technology Beijing. Herresearch interests include powder metallurgy titanium andtitanium alloys, aluminum and aluminum alloys, copperand copper alloys, 3D printing, iron-based alloys, self-propa-gating high temperature synthesis (SHS), and magneticmaterials.
Qian Qin is Ph.D. candidate in Engineering at the Universityof Science and Technology Beijing. He is engaged in theresearches on powder metallurgy iron-based materials.
Biao Zhang is a master in Engineering graduated from theUniversity of Science and Technology Beijing. He is engagedin the researches on powder metallurgy iron-based materials.
Alex A. Volinsky is an Associate Professor at the Universityof South Florida, USA. He is an expert in thin films proces-sing, mechanical properties and characterization, adhesionand fracture of thin films, nanoindentation, pattern for-mation, irradiated materials properties and X-Raydiffraction.
Zhimeng Guo is a Professor and Ph.D. supervisor working atthe University of Science and Technology Beijing. He offerscourses in Powder Metallurgy. He is an expert in powdermetallurgy titanium and titanium alloys, aluminum andaluminum alloys, copper and copper alloys, 3D printing,iron-based alloys, dispersion strengthened materials, radiofrequency inductively coupled plasma spheroidization tech-nology, self-propagating high temperature synthesis (SHS),advanced powder metallurgy technologies and materials.
ORCID
Alex A. Volinsky http://orcid.org/0000-0002-8520-6248
References
[1] Mesquita RA, Barbosa CA. High-speed steels pro-duced by conventional casting, spray forming andpowder metallurgy. Mater Sci Forum. 2005;498–499:244–250.
[2] Sustarsic B, Kosec L, Jenko M, et al. Vacuum sinteringof water-atomised HSS powders with MoS2 additions.Vacuum. 2001;61:471–477.
[3] Sustarsic B, Kosec L, Kosec M, et al. The influence ofMoS2 additions on the densification of water-atomizedHSS powders. J Mater Process Technol. 2006;173:291–300.
[4] Škapin SD, Jenko M, Leskovšek V, et al.Characterization of the carbides and the martensitephase in powder-metallurgy high-speed steel. MaterCharact. 2010;61:452–458.
[5] Zhang X, Liu W, Sun D, et al. The transformation ofcarbides during austenization and its effect on thewear resistance of high speed steel rolls. Metall MaterTrans A. 2007;38:499–505.
[6] Ding P, Shi G, Zhou S. As-cast carbides in high-speedsteels. Metall Mater Trans A. 1993;24:1265–1272.
[7] Zhang L, Zhang X, Li L, et al. Evolution of the micro-structure and mechanical properties of powder metal-lurgical high-speed steel S390 after heat treatment. JAlloys Compd. 2017;740:766–773.
[8] Grinder O. The HIP way to make cleaner, better steels.Met Powder Rep. 2007;62:16-18.
[9] Wright CS, Ogel B. Supersolidus sintering of highspeed steels: part 1: sintering of molybdenum basedalloys. Powder Metall. 1993;36:213–219.
[10] Wright CS, Ogel B, Lemoisson F, et al. Supersolidussintering of high speed steels: part 2: sintering oftungsten based alloys. Powder Metall. 1995;38:221–229.
[11] Zhang D, Li Z, Xie L, et al. Powder metallurgy of highspeed-steel produced by solid state sintering and heattreatment. Int J Mater Res. 2015;106:870–876.
[12] Zhang QK, Jiang Y, ShenWJ, et al. Direct fabrication ofhigh-performance high speed steel products enhancedby LaB6. Mater Des. 2016;112:469–478.
[13] Giménez S, Zubizarreta C, Trabadelo V, et al. Sinteringbehaviour and microstructure development of T42powder metallurgy high speed steel under differentprocessing conditions. Mater Sci Eng A.2008;480:130–137.
[14] Aguirre I, Gimenez S, Talacchia S, et al. Effect of nitro-gen on supersolidus sintering of modified M35M highspeed steel. Powder Metall. 1999;42:353–357.
[15] Bolton JD, Gant AJ. Fracture in ceramic-reinforcedmetal matrix composites based on high-speed steel. JMater Sci. 1998;33(4):939–953.
[16] Bolton JD, Gant AJ. Microstructural development andsintering kinetics in ceramic reinforced high speed steelmetal matrix composites. Powder Metall. 1997;40(2):143–151.
[17] Gordo E, Velasco F, Antón N, et al. Wear mechanismsin high speed steel reinforced with (NbC)p and (TaC)pMMCs. Wear. 2000;239(2):251–259.
[18] Pines ML, Bruck HA. Pressureless sintering of particle-reinforced metal-ceramic composites for functionallygraded materials: part I. Porosity reduction models.Acta Mater. 2006;54:1457–1465.
[20] Kim HC, Oh DY, Shon IJ. Sintering of nanophase WC-15vol.%Co hard metals by rapid sintering process. Int JRefract Met Hard Mater. 2004;22:197–203.
[21] Trabadelo V, Giménez S, Iturriza I. Microstructuralcharacterisation of vacuum sintered T42 powdermetallurgy high-speed steel after heat treatments.Mater Sci Eng A. 2009;499:360–367.
[22] Godec M, Večko Pirtovšek T, Šetina Batič B, et al.Surface and bulk carbide transformations in high-speed steel. Sci Rep. 2015;5:1–11.
[23] Gimenez S, Iturriza I. Microstructural characterisationof powder metallurgy M35MHV HSS as a function ofthe processing route. J Mater Process Technol.2003;143–144:555–560.
[24] Xu HX, Zhu WL, Jiang JQ, et al. Refining carbidedimensions in AISI M2 high speed steel by increasingsolidification rates and spheroidising heat treatment. JMater Sci Technol. 2013;30:116–122.
[25] Hetzner DW. Refining carbide size distributions in M1high speed steel by processing and alloying. MaterCharact. 2001;46:175–182.
