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Microstructure evolution during hot rolling and heat treatment of the spray formed Vanadis 4 cold work steel
M A TE RI A L S C H A RAC TE RI ZA T ION 5 9 ( 2 00 8 ) 1 0 0 7–1 0 1 4
Microstructure evolution during hot rolling and heat treatment of the spray formed Vanadis 4 cold work steel Fei Yan a,⁎, Haisheng Shi b , Bingzhong Jin b , Junfei Fan b , Zhou Xu a a
Key Laboratory for High Temperature Materials and Tests of Ministry of Education, Shanghai Jiao Tong University, 1954 Hua Shan Road, Shanghai 200030, China b Shanghai Baosteel Research Institute, Shanghai 201900, China
AR TIC LE D ATA
ABSTR ACT
Article history:
A high alloyed Vanadis 4 steel was produced by spray forming, and the microstructure
Received 25 March 2007
evolution during hot rolling and annealing processes was characterized. It was found that
Received in revised form
the as-sprayed Vanadis 4 steel has a homogeneous and fine microstructure with uniform
3 August 2007
dispersion of carbides. The hot rolling temperature is the key factor in controlling the
Accepted 13 August 2007
evolution of type, morphology and distribution of carbides, as well as the matrix microstructure of the as-rolled steels. The optimized processing parameters for the as-
Keywords:
sprayed Vanadis 4 steel are rolling at 1050 °C and annealing at 900 °C. The microstructural
Spray forming
evolution mechanisms during hot rolling and annealing are determined according to the
Vanadis 4 steel
microstructural analysis of the material at different stages.
Microstructure evolution
1.
Introduction
Vanadis 4 (V4) steel is a powder metallurgical (PM) high alloyed cold work tool steel containing (wt.%) 1.5 C, 1.0 Si, 0.4 Mn, 8.0 Cr, 1.5 Mo, 4.0 V. The steel offers an extremely good combination of wear resistance and ductility for high performance tools. However, high cost and complicated working procedures are the main disadvantages of powder metallurgy, which restrains the wide application of the steel. For some near-net-shape applications, spray forming is expected to replace powder metallurgy to produce high alloyed steels. Spray forming, also called spray casting or spray deposition, is the inert gas atomization of a liquid metal stream into various sized droplets which are then propelled away from the region of atomization by the fast flowing, atomizing gas [1]. The droplets were collected and solidified on a substrate, and finally deposit into a coherent, near fully dense product. Rapid solidification effects inherent in the spray forming process due to high heat exchange rate at the droplet–gas interface and also on the deposition surface ensures considerable chemical and microstructural homogeneity of the product
© 2007 Elsevier Inc. All rights reserved.
[2]. As a result, the material microstructures by spray forming differ significantly from those of materials by both conventional casting and powder metallurgy. Application of the spray forming technique in the field of steels has been focused on high alloyed steels where the benefits of the technique especially the possibility to produce a fine-grained, segregation-free microstructure are very promising [3–6]. The cost benefit derives from the single-step operation of converting molten alloy directly into a semi-finished product. Therefore, spray forming has emerged as a key competitor for existing technologies, especially powder metallurgy. But, as reported in many papers, an excess solid fraction in the spray generates a product with some pores due to insufficient liquid phase available to provide bonding of particles during solidification of the spray formed materials [6,9,10]. After spray forming, the billet need to be forged or hot rolled to eliminate the remaining porosity. Although many efforts have been made by using spray forming to produce high alloyed steels [3–8], there is little work published on the microstructural evolution behavior during the whole hot working process [4]. Furthermore, to our knowledge, few works have been published on comparing
⁎ Corresponding author. Tel.: +86 21 26649626; fax: +86 21 26643987. E-mail address: [email protected] (F. Yan). 1044-5803/$ – see front matter © 2007 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2007.08.012
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the microstructure of spray formed steel with commercial powder metallurgical equivalent. The V4 cold work steel was produced by spray forming and its microstructure was studied. Detailed study was focused on the microstructural evolution during hot rolling and annealing in order to optimize processing conditions and to examine their effects on microstructure and carbides. The objectives of the present work were to: (1) select the optimum parameters of hot rolling and annealing for the as-sprayed V4 steel; (2) investigate the microstructural evolution in this steel during hot working and heat treatment process; and (3) compare the microstructure of the as-sprayed steel after hot rolling and annealing with that of a commercial powder metallurgical equivalent manufactured by UDDEHOLM, Sweden.
2.
Experimental Procedure
The steels were melted in a vacuum induction furnace and then cast into rods as the feedstock for spray forming. The feedstock was heated in an induction-furnace and soaked at above the melting point for 20 min. The molten metal flow rate was set at approximately 0.1 kg/s, using N2 with a pressure of 2.2 MPa as the atomizing gas. The atomized droplets were cooled and driven towards a revolving substrate to form a dense product. The distance from the nozzle to the substrate was set as 360 mm and the copper substrate was rotated at a speed of 10 rpm during atomization and deposition. The spray-forming process was completed in about 40 s, and a gauss-shaped billet with about 130 mm in diameter and 30 mm in height was obtained. Hot rolling and annealing were carried out based on the critical temperatures, which were determined by dilatometry using a “ThermecMastor_Z” thermomechanical simulator. Cylindrical specimens 12 mm in height and 8 mm in diameter were heated to 1100 °C at a rate of 0.2 °C/s, being held for 10 min, and then cooled firstly to 600 °C at a rate of 0.1 °C/s and finally to 200 °C at a rate of 1 °C/s. Phase transformations were detected by monitoring the fractional change in dilatation with temperature. The spray formed billet was machined into specimens with a thickness of 16 mm. The specimens were then heated at a rate of 10 °C/min to the rolling temperatures of 850 °C, 900 °C, 950 °C, 1050 °C, 1150 °C, respectively. After holding for 15 min, some specimens were quenched in water in order to hold the carbides' morphology before hot rolling. Other specimens were rolled in a single 60% reduction pass, and then cooled to room temperature in sand. Three groups of as-rolled specimens were annealed at 850 °C, 900 °C, and 950 °C respectively, with each group containing five specimens rolled at different temperatures. After being isothermally held for 2 h, the specimens were cooled to 500 °C at a rate of 30 °C/h, and then cooled in air. Microstructural observations were made optically, with an S-4200 field emission scanning electron microscope (SEM) and with an H-800 transmission electron microscope (TEM). Thin foils for transmission electron microscopic studies were prepared from 3 mm disks, grounded to a thickness of about 50 μm and electropolished in an electrolyte containing 5% perchloric and 95% ethanol at −20 °C. Bright field and dark
field imaging and selected area diffraction patterns were used to identify the carbides. The specimens for Vicker's hardness testing were performed on a Vickers hardness tester (Model: FR-3E) with 30 kg load and each measurement was tested for 8 s. The average hardness of five measurements was obtained as the final hardness.
3.
Results and Discussion
3.1. Microstructure of the AS-sprayed V4 Steel and Phase Transformation Temperature An optical micrograph of the as-sprayed microstructure is shown in Fig. 1a. Spray forming resulted in a substantial reduction in microstructural scale for both grain size and carbide size compared with conventionally cast equivalent [11]. Fine, uniformly distributed equiaxed grains in diameters ranging from 8 to 10 μm are observed. In addition, large primary and interconnected eutectic carbides observed in the conventional cast material are replaced by a more uniformly distributed spheroidal carbides in the as-sprayed material. As shown in Fig. 1b, observation by TEM shows that the matrix of the microstructure of the as-sprayed material was predominantly twinned martensite. Selected area diffraction patterns confirmed the presence of V-rich MC and Cr-base M7C3 type
Fig. 1 – (a) Optical micrograph of the as-sprayed V4 steel, (b) TEM morphology of the twined matrensite.
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teresting to find that an obvious flexure appears during heating between temperatures ranging from 650 °C to 760 °C, indicating the decomposition of martensite. The cooling curve indicates that the high alloyed austenite has a strong stability, as no obvious transformations could be observed other than the bainite reaction which occurred only at very slow cooling rates.
3.2.
Fig. 2 – Measured dilatation variation with temperature.
carbides. Most of the MC and some big M7C3 carbides ranging from 0.5 to 2 μm in diameter are distributed on the grain boundaries. Other small M7C3 carbides with the diameter of about 180 nm distribute in the grains. The variation in specimen dimensions with temperature during heating and cooling process by the dilatometry experiment is shown in Fig. 2. The transformation temperatures, 845 °C (Ac1s) and 890 °C (Ac1f), are defined as the points where the curve leaves the tangent to the linear portion. It is in-
Microstructure Evolution During Hot Rolling
Fig. 3 shows the SEM micrographs of the microstructures before hot rolling at different temperatures. It can be seen that precipitation occurred in all of the cases and obvious differences lie in the amount and morphology of the precipitates. Fig. 3 indicates that the amount of the small carbides decreases with increasing temperature. As shown in Fig. 3a, a larger number of small carbides appear at 850 °C. TEM observation shows that they are M3C or M7C3, however, the bigger, irregular phases, as shown by arrows, are MC (VC) carbides. When the temperature is elevated to 950 °C or even higher, only MC and M7C3 can be found. Thus, reheating temperature leads to major microstructural changes involving phase transformation, dissolution and spheroidization of carbides formed in the as-sprayed steels. Fig. 4 shows the SEM observation on the microstructures of the as-rolled steels which were rolled at 850 °C, 950 °C, 1050 °C, and 1150 °C, respectively. The stable MC carbide particles
Fig. 3 – SEM micrographs of the as-sprayed steels quenched at: (a) 850 °C; (b) 950 °C; (c) 1050 °C and (d) 1150 °C.
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Fig. 4 – SEM micrographs of the as-sprayed steel rolled at: (a) 850 °C; (b) 950 °C; (c) 1050 °C and (d) 1150 °C.
show little tendency to coarsen. The M7C3 carbide, unlike MC, is very sensitive to the rolling temperature, i.e., both the size and morphology of the M7C3 carbides vary with temperature. For example, TEM observation suggests that many coarse carbides shown in Fig. 4a are of M7C3 type. But their size and amount are all decreased with increased rolling temperature. Fig. 5a is the TEM image of those small carbides shown in Fig. 4a. It was confirmed by electron diffraction patterns that most of the small carbides are of M7C3 type, and some M3C carbides are also found. Electron diffraction patterns of M3C and M7C3 carbides are shown in Fig. 5b and c, respectively. The characteristics shown in Fig. 3a and Fig. 4a suggest that the microstructural evolution process at 850 °C differs significantly from those at other temperatures. This can be attributed to the difference of the diffusion coefficient of the alloy elements. Nucleation accompanied by growth and followed by coarsening constitutes the processes during the hot rolling and subsequent cooling process. The nucleation rate as the derivative of the precipitate density N vs time t can be expressed as [12]: s dN DG⁎ 4kR⁎2 DXC0 ¼ N0 Zb⁎ exp and b⁎ ¼ exp : dt t kt a4
ð1Þ
Where N0 is the number of nucleation sites per unit volume, Z is the Zeldovich factor, τ is the incubation time, ΔG⁎ is the nucleation driving force, R⁎ is the nucleation radius,
D is the diffusion coefficient of solute atoms in the matrix, and XC0 is the initial solute mole fraction. The coarsening of second phase particles as a result of a reduction in interfacial energy was theoretically treated by Lifshitz and Slyozov [13] and Wagner where the result was expressed as: ¯r 3 ¯r 30 ¼
2 8gDce Vm t: 9RT
ð2Þ
Where r̄ is the average particle radius at time t, r̄0 is the average particle radius at the onset of coarsening, γ is the interfacial free-energy of the particle–matrix interface, D is the diffusion coefficient of solute atoms in the matrix, ce is the concentration of solute in the matrix in equilibrium with a particle of infinite size, Vm is the molar volume of the particle, R is the gas constant, and T is the temperature. The nucleation rate and coarsening of the carbides are all governed by the diffusion coefficient D of solute atoms, and therefore, the magnitude of the diffusion coefficients provides a useful comparison basis. The diffusion coefficient D can be presented in terms of an Arrhenius equation: Q : D ¼ D0 exp RT
ð3Þ
Where D0 is the pre-experimental factor for the diffusion coefficient, and Q is the activation energy for diffusion.
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to the tempering of martensite around 700 °C [15]. In the early stages of the decomposition, the reaction is controlled by carbon diffusion and is thus very rapid. The change of the material's volume occurs because a large number of M3C carbides precipitate from matrix. This is why an obvious flexure appears in dilatability–temperature curve (Fig. 2) between the temperatures of 650 °C and 760 °C. The amount of the precipitates below the A1 temperature was tremendous, as shown in Fig. 3a. A metastable supersaturated phase will not only generate nuclei but also cause deposition on the generated nuclei and consequent growth of the clusters. This decreases the supersaturation, and correspondingly increases the critical nucleus size, and therefore promotes the coarsening process [16]. As the result of this, the Gibbs–Thomson effect becomes so significant that smaller clusters dissolve, transfer their mass to larger growing clusters and coarsen the size distribution [17]. When small clusters shrink to their critical nucleus size, they become thermodynamically unstable, and spontaneously disintegrate and eventually vanish [18]. This is why a large number of clusters nucleate and then grow, Fig. 3a, and many coarse carbides appear, Fig. 4a. Occasionally coarse carbides, which are favorably situated at the grain corners, appear to have grown from those small carbides close to them. This conclusion is supported by the observation that such particles are frequently polygonal in shape, and there are comparatively fewer small carbides that
Fig. 5 – (a) TEM micrograph of the carbides in the material rolled at 850 °C, electron diffraction patterns in [100] zone axis of M3C particle (b), and in [431] zone of M7C3 particle (c).
According to Eq. (3), the diffusion coefficients of carbon (with D0 = 0.738 cm2 · s− 1 and Q = 158.98 kJ · mol− 1 [14]) and chromium (with D0 = 4.08 cm2 · s− 1 and Q = 286.8 kJ · mol− 1 [14]) at 850 °C are 2.97 × 10− 8 cm2 · s− 1 and 1.86 × 10− 13cm2 · s− 1, respectively. The diffusion coefficients between carbon and chromium are of great difference. Since martensite is not an equilibrium phase, when steel is heated below the eutectoid temperature, the thermodynamically stable phases α and M3C start to precipitate. At temperatures ranging from 650 °C to 760 °C, precipitation of M3C carbides can be written as: α′ ⇒ α + M3C. This reaction is similar
Fig. 6 – (a) TEM micrograph of the polygonal carbide, (b) its electron diffraction pattern in [312] zone axis.
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exist in the vicinity of the coarse particles. On the other hand, in many cases, the firstly formed carbides are non-equilibrium phase and they are replaced by more stable forms gradually in the precipitation sequence [19]. After the precipitation of M3C carbides, the kinetics of the diffusion mechanism changes from carbon diffusion control to chromium diffusion control. Therefore, the observed small M7C3 carbides, as shown in Fig. 5, may be attributed to the transformation from M3C to M7C3 carbides. The possibility of M7C3 carbide formation from M3C carbides was also proved by Homolová [20] and Liu [21,22]. During the heat up process, the carbide precipitation and dissolution is a competition process. At relatively lower temperature the precipitation dominates the process while at high temperature dissolution dominates. Therefore, the decrease of the amount of the small carbides can be attributed to the dissolution with the increase of temperature, and the saturation of the alloy elements in the matrix will increase also. The microstructure of the steel which rolled at 1050 °C shows the best uniform distribution of the carbides with ideal average size, Fig. 4c. Some big polygonal carbides appear at the triplejunctions of grain boundaries in the material which was rolled at 1150 °C, Fig. 4d. Those polygonal carbides were not found in the materials rolled at other temperatures. TEM morphology of this kind of carbide and its corresponding electron diffraction pattern indicate that this carbide is of M7C3, Fig 6. After hot rolling, the supersaturated alloying elements will precipitate during cooling. The precipitation of the carbides is determined by two basic factors. One is that the alloying elements in the matrix should be in supersaturation state, and the other that the atoms have enough ability to diffuse. Both of these factors closely correlate with the temperature. It is thus reasonable to infer that the supersaturation and the diffusion abilities of the alloy atoms are at their highest at 1150 °C in comparison with those at other temperatures. In many diffusion phase transformation and precipitation process, nucleation of the product phase occurs heterogeneously at some preferential nucleation sites in the matrix such as the grain boundaries, dislocations and second phases. Grain boundaries often contain a high density of defects, which can be increased by plastic deformation [23]. Therefore,
Fig. 7 – Hardness of the as-sprayed steels rolled at different temperatures.
Fig. 8 – TEM micrograph of martensite in the as-rolled steel.
carbide precipitation on grain boundaries will be more easily, and precipitates which are growing on boundaries will gain a bigger proportion of their material from within the grains. Although the dissolution of chromium carbides in matrix has been reported by researchers as a more rapid process than the dissolution of other carbides, the complete dissolution of M7C3 carbides is very difficult even at very high austenitizing temperature [24]. When the as-sprayed material was heated to 1150 °C, many M7C3 carbides on the grain boundary still not dissolved completely. They act as preferential nucleation sites and grow up quickly. Thus, big polygonal M7C3 carbides formed and the various shapes of grain boundary M7C3 carbides may possibly be explained by variation of high angle boundaries. There is evidence from Fig. 4d that the most frequently occurring form of grain boundary precipitate is discontinuous. The limited boundary movement observed can be explained in the manner that the precipitation of the carbides in the grain boundaries act as effective obstacles of the growth of the grains, which ensures fine grains remained in elevated temperature. Results of the hardness as a function of the rolling temperature are shown in Fig. 7. It can be seen that the hardness increases with increased rolling temperature. Especially there is a sharp increase from 900 °C to 950 °C. The origination of this effect can be ascribed to the completely different matrix microstructure after hot rolling. TEM observation shows the matrix of the as rolled steels rolled at 850 °C and 900 °C consists of ferrite, as shown in Fig. 5a. Whereas the matrix of the as rolled steels which rolled at or higher than 950 °C consists of martensite, and the representative morphology of the matrix is shown in Fig. 8 (by the steel rolled at 1050 °C). The as-rolled microstructure was controlled by the phase transformation during hot rolling and cooling process, and the rolling temperature is actually a key in this process. When the steel is heated to 900 °C (slightly above A1), little carbide dissolves and the contents of both carbon and alloying elements in austenite are lower. The unstable austenite could transform into ferrite during cooling. But when the temperature is elevated to 950 °C or higher, more alloy elements dissolve in the austenite and improve the hardenability of the steel. Martensite can be therefore form even
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when the steel is cooled in sand. As a result, the completely different microstructure, ferrite and martensite, cause the sharp increase of the hardness from 900 °C to 950 °C.
3.3.
Microstructure Evolution During Annealing
Table 1 – Hardness of the as-rolled steels annealed at different temperatures (HV) Rolling temperature (°C)
Annealing temperature 850 °C
900 °C
950 °C
288 268 278 289 280
243 240 237 246 240
231 230 228 238 219
850 °C
All the as-rolled steels were annealed at different temperatures in order to find the optimum combination of hot rolling and annealing parameters. The best microstructure, with ideal average size and uniform distribution of the carbides, is shown in Fig. 9a. It was obtained by the as-sprayed steel rolled at 1050 °C and then annealed at 900 °C, whereas all other asannealed steels showed inhomogeneous distribution of carbide size to some extent. It can be seen from Fig. 9 that the average carbide size in the spray formed steel is even smaller than the equivalent in the commercial PM V4 steel. TEM observation shows that the as-annealed steel consists of ferrite and homogeneously distributed granular MC and M7C3 carbides. Table 1 compares Vickers hardness of all the asannealed steels. The hardness of those steels which annealed at 900 °C is close to that of the commercial PM V4 steel (HV: 243). The results give strong evidence that the high alloyed V4 steel can be produced by spray forming. Annealing of the as-rolled steels modifies the microstructure in three ways: (1) recrystallization of the matrix; (2) transform the matrix into soft ferrite; and (3) coarsening and further spheroidization of carbides. Microstructural
Fig. 9 – Comparison of V4 annealed microstructures: (a) spray forming, (b) powder metallurgy.
900 °C 950 °C 1050 °C 1150 °C
evolution processes during annealing varied according to the variations of the as rolled steels' matrix. As to those materials, which were rolled at or higher than 950 °C, their matrix is composed of martensite. Similar to the reheating process of the as-sprayed material, when they were heated to elevated temperature the martensite will decompose into ferrite and M3C carbides quickly. But this does not happen during the reheating process to the steels rolled at 850 °C and 900 °C as their matrix is ferrite. In chromium steel the M3C carbides usually have stronger tendency to dissolve in matrix than M7C3 carbides because the later is thermodynamically more stable [24]. Upon heating up to austenitizing temperature, there is a likelihood of in-situ transformation from undissolved M3C carbides to M7C3 carbides. The transformation may begin with the dissolution of M3C followed by in situ nucleation and growth of M7C3 carbides [24], or may proceed by the eutectoid reaction of M3C ⇒ γ + M7C3 [22]. It is well known that a distribution of immobile particles in a solid matrix tends to lower its interfacial free energy by transport of matter from the small to the larger members, thereby diminishing the total particles number but increasing the average particle size. This coarsening process also is known as Ostwald ripening [25,26]. As to the MC and M7C3 carbides, V-rich MC carbides have little tendency to coarsen but Cr-base M7C3 carbides are prone to grow up at elevated temperature. In addition, the growth of dispersed particles in a saturated solution by a diffusion mechanism depends on many variables. As has reported by Greenwood [27], the main factors promoting stability (slow growth-rate) are particles of large mean radius, high density, and low molecular weight, with low values of the solute diffusion coefficient. The carbides' size and distribution of the as-annealed steels are influenced by the microstructures of the as-rolled steels. In general, attractive annealed microstructure can be obtained from the as-rolled steels with fine, homogeneously distributed carbides. Two representative as-annealed steels, which were obtained by rolling at 850 °C and 1050 °C respectively and then all annealed at 900 °C, could be taken to explain the difference. As shown in Fig. 4a, the radii of the M7C3 carbides varied greatly. During the isothermal holding process during annealing, the crucial effect of interfacial curvature on energy, and hence on particle size (Gibbs–Thomson effect), causes larger M7C3 carbides to grow at the expense of the smaller particles, so that smaller particles dissolve and eventually vanish. Thus, it is hard to obtain ideal microstructure by the initial
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incongruous size and distribution of the carbides. On the contrary, as to the specimens rolled at 1050 °C, the carbides are highly dispersed and have a uniform size distribution. Then they have higher stability and grow slowly according to Greenwood's conclusion [27]. No polarizing coarsening behavior as discussed above will appear and it is more likely to obtain ideal equal carbides' size and distribution.
4.
Conclusions
1. Spray forming resulted in a significant refinement of both grain size and carbide size of the high alloyed V4 cold work steel. The optimized processing parameters for the assprayed V4 steel are rolling at 1050 °C and then annealing at 900 °C. No interconnected eutectic carbide structure was observed and fine spheroidal MC and M7C3 carbides were uniformly distributed in the matrix. 2. The hot rolling temperature is the key factor in controlling the evolution of type, morphology and distribution of carbides, as well as in controlling the matrix microstructure of the as-rolled steel. The matrix consists of ferrite when the materials rolled at 850 °C and 900 °C. But when the steels were rolled at or higher than 950 °C, martensite will be obtained because of enhanced hardenability by dissolving of alloy elements in austenite. 3. M3C is the only transient carbide precipitate below the A1 temperature during the reheating of the as-sprayed steel. Irregular M7C3 carbides are preferentially precipitated at the grain boundaries during cooling after the steel was rolled at 1150 °C. 4. The size and distribution of carbides in the as-annealed steels are influenced by the microstructures of the as-rolled steels. Attractive annealed microstructure can be obtained from the as-rolled steels with fine, homogeneously distributed carbides.
Acknowledgements This research was financially supported by Shanghai Baosteel Technology Center and Key Laboratory for High Temperature Materials and Tests of Ministry of Education, Shanghai Jiao Tong University. The author would like to thank Dr Wei Wang for many useful suggestions and comments.
REFERENCES [1] Grant PS. Spray forming. Prog Mater Sci 1995;39:497–545. [2] Ojha SN, Jha JN, Singh SN. Microstructural modifications in Al–Si eutectic alloy produced by spray deposition. Scr Metall Mater 1991;25:443–7. [3] Zhang JG, Xu HB, Shi HS, Wu JS, Sun DS. Microstructure and properties of spray formed Cr12MoV steel for rolls. J Mater Process Technol 2001;111:79–84. [4] Rodenburg C, Krzyzanowski M, Beynon JH, Rainforth WM. Hot workability of spray-fromed AISI M3:2 high-speed steel. Mater Sci Eng A 2004;386:420–7.
[5] Spangel S, Matthaei-Schulz E, Schulz A, Vetters H, Mayr P. Influence of carbon and chromium content and preform shape on the microstructure of spray formed steel deposits. Mater Sci Eng A 2002;326:26–39. [6] Mesquita RA, Barbosa CA. Spray forming high speed steel — properties and processing. Mater Sci Eng A 2004;383:87–95. [7] Hanlon DH, Rainforth WM. The rolling sliding wear response of conventionally proceed and spray formed high speed steel at ambient and elevated temperature. Wear 2003;255:956–66. [8] Yang YF, Hannula SP. Sprayformed hot work steels for rapid tooling. J Mater Sci Technol 2003;19:169–72. [9] Upadhyaya A, Mishra NS, Ojha SN. Microstructural control by spray forming and wear characteristics of a Babbit alloy. J Mater Sci 1997;32:3227–35. [10] Dobrzański LA. Effects of chemical composition and processing conditions on the structure and properties of high-speed steels. J Mater Process Technol 1995;48:727–37. [11] Yan F, et al. Microstructure of the spray formed Vanadis 4 steel and its ultrafine structure. Mater Charact 2007, doi:10.1016/j.matchar.2007.04.019. [12] Michel P, Alexis D. Microscopic modeling of simultaneous two-phase precipitation application to carbide precipitation in low-carbon steels. Mater Sci Eng A 2003;360:214–9. [13] Lifshitz IM, Slyozov VV. The kinetics of precipitation from supersaturated solid solutions. J Phys Chen Solids 1961;19:35–50. [14] Yu YL. Metallic theories. Bei Jing: Metallurgical Industry Press; 2003. p. P213–4. [In Chinese]. [15] Shtansky DV, Inden G. Phase transformation in Fe–Mo–C and Fe–W–C steels—I. The structural evolution during tempering at 700 °C. Acta Mater 1997;45:2861–78. [16] McClurg RB, Flagan RC. Critical comparison of droplet models in homogeneous nucleation theory. J Colloid Interface Sci 1998;201:194–9. [17] Madras G, Mccoy BJ. Temperature effects on the transition from nucleation and growth to Ostwald ripening. Chem Eng Sci 2004;59:2753–65. [18] Madras G, Mccoy BJ. Distribution kinetics theory of Ostwald ripening. J Chem Phys 2001;115:6699–706. [19] Kaneko K, Matsumura S, Sadakata A, Fujita K, Moon WJ, Ozaki S, et al. Characterization of carbides at different boundaries of 9Cr-steel. Mater Sci Eng A 2004;374:82–9. [20] Homolová V, Janovec J, Kroupa A. Experimental and thermodynamic studies of phase transformations in Cr–V low alloy steels. Mater Sci Eng A 2002;335:290–7. [21] Liu ZK, Höglund L, Jönsson B, Ågren J. An experimental and theoretical study of cementite dissolution in an Fe–Cr–C alloy. Metall Trans A 1991;22:1745–52. [22] Liu ZK, Ågren J. Morphology of cementite decomposition in an Fe–Cr–C alloy. Metall Trans A 1991;22:1753–9. [23] Adamson JP, Martin JW. The nucleation of M23C6 carbide particles in the grain boundaries of an austenitic stainless steel. Acta Metall 1971;19:1015–8. [24] Wu C, Sahajwalla V, Krauklis P. The effect of austenitizing process on the hardening behaviour of Cr–Mo–Mn–C air-hardening cast tool steel. ISIJ Inter 1996;36:347–53. [25] Speight MV. Growth kinetics of grain-boundary precipitates. Acta Metall 1968;16:133–5. [26] Ardell AJ. On the coarsening of grain boundary precipitates. Acta Metall 1972;20:601–9. [27] Greenwood GW. The growth of dispersed precipitates in solutions. Acta Metall 1956;4:243–8.
Microstructure evolution during hot rolling and heat treatment of the spray formed Vanadis 4 cold work steel Fei Yan a,⁎, Haisheng Shi b , Bingzhong Jin b , Junfei Fan b , Zhou Xu a a
Key Laboratory for High Temperature Materials and Tests of Ministry of Education, Shanghai Jiao Tong University, 1954 Hua Shan Road, Shanghai 200030, China b Shanghai Baosteel Research Institute, Shanghai 201900, China
AR TIC LE D ATA
ABSTR ACT
Article history:
A high alloyed Vanadis 4 steel was produced by spray forming, and the microstructure
Received 25 March 2007
evolution during hot rolling and annealing processes was characterized. It was found that
Received in revised form
the as-sprayed Vanadis 4 steel has a homogeneous and fine microstructure with uniform
3 August 2007
dispersion of carbides. The hot rolling temperature is the key factor in controlling the
Accepted 13 August 2007
evolution of type, morphology and distribution of carbides, as well as the matrix microstructure of the as-rolled steels. The optimized processing parameters for the as-
Keywords:
sprayed Vanadis 4 steel are rolling at 1050 °C and annealing at 900 °C. The microstructural
Spray forming
evolution mechanisms during hot rolling and annealing are determined according to the
Vanadis 4 steel
microstructural analysis of the material at different stages.
Microstructure evolution
1.
Introduction
Vanadis 4 (V4) steel is a powder metallurgical (PM) high alloyed cold work tool steel containing (wt.%) 1.5 C, 1.0 Si, 0.4 Mn, 8.0 Cr, 1.5 Mo, 4.0 V. The steel offers an extremely good combination of wear resistance and ductility for high performance tools. However, high cost and complicated working procedures are the main disadvantages of powder metallurgy, which restrains the wide application of the steel. For some near-net-shape applications, spray forming is expected to replace powder metallurgy to produce high alloyed steels. Spray forming, also called spray casting or spray deposition, is the inert gas atomization of a liquid metal stream into various sized droplets which are then propelled away from the region of atomization by the fast flowing, atomizing gas [1]. The droplets were collected and solidified on a substrate, and finally deposit into a coherent, near fully dense product. Rapid solidification effects inherent in the spray forming process due to high heat exchange rate at the droplet–gas interface and also on the deposition surface ensures considerable chemical and microstructural homogeneity of the product
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[2]. As a result, the material microstructures by spray forming differ significantly from those of materials by both conventional casting and powder metallurgy. Application of the spray forming technique in the field of steels has been focused on high alloyed steels where the benefits of the technique especially the possibility to produce a fine-grained, segregation-free microstructure are very promising [3–6]. The cost benefit derives from the single-step operation of converting molten alloy directly into a semi-finished product. Therefore, spray forming has emerged as a key competitor for existing technologies, especially powder metallurgy. But, as reported in many papers, an excess solid fraction in the spray generates a product with some pores due to insufficient liquid phase available to provide bonding of particles during solidification of the spray formed materials [6,9,10]. After spray forming, the billet need to be forged or hot rolled to eliminate the remaining porosity. Although many efforts have been made by using spray forming to produce high alloyed steels [3–8], there is little work published on the microstructural evolution behavior during the whole hot working process [4]. Furthermore, to our knowledge, few works have been published on comparing
⁎ Corresponding author. Tel.: +86 21 26649626; fax: +86 21 26643987. E-mail address: [email protected] (F. Yan). 1044-5803/$ – see front matter © 2007 Elsevier Inc. All rights reserved. doi:10.1016/j.matchar.2007.08.012
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the microstructure of spray formed steel with commercial powder metallurgical equivalent. The V4 cold work steel was produced by spray forming and its microstructure was studied. Detailed study was focused on the microstructural evolution during hot rolling and annealing in order to optimize processing conditions and to examine their effects on microstructure and carbides. The objectives of the present work were to: (1) select the optimum parameters of hot rolling and annealing for the as-sprayed V4 steel; (2) investigate the microstructural evolution in this steel during hot working and heat treatment process; and (3) compare the microstructure of the as-sprayed steel after hot rolling and annealing with that of a commercial powder metallurgical equivalent manufactured by UDDEHOLM, Sweden.
2.
Experimental Procedure
The steels were melted in a vacuum induction furnace and then cast into rods as the feedstock for spray forming. The feedstock was heated in an induction-furnace and soaked at above the melting point for 20 min. The molten metal flow rate was set at approximately 0.1 kg/s, using N2 with a pressure of 2.2 MPa as the atomizing gas. The atomized droplets were cooled and driven towards a revolving substrate to form a dense product. The distance from the nozzle to the substrate was set as 360 mm and the copper substrate was rotated at a speed of 10 rpm during atomization and deposition. The spray-forming process was completed in about 40 s, and a gauss-shaped billet with about 130 mm in diameter and 30 mm in height was obtained. Hot rolling and annealing were carried out based on the critical temperatures, which were determined by dilatometry using a “ThermecMastor_Z” thermomechanical simulator. Cylindrical specimens 12 mm in height and 8 mm in diameter were heated to 1100 °C at a rate of 0.2 °C/s, being held for 10 min, and then cooled firstly to 600 °C at a rate of 0.1 °C/s and finally to 200 °C at a rate of 1 °C/s. Phase transformations were detected by monitoring the fractional change in dilatation with temperature. The spray formed billet was machined into specimens with a thickness of 16 mm. The specimens were then heated at a rate of 10 °C/min to the rolling temperatures of 850 °C, 900 °C, 950 °C, 1050 °C, 1150 °C, respectively. After holding for 15 min, some specimens were quenched in water in order to hold the carbides' morphology before hot rolling. Other specimens were rolled in a single 60% reduction pass, and then cooled to room temperature in sand. Three groups of as-rolled specimens were annealed at 850 °C, 900 °C, and 950 °C respectively, with each group containing five specimens rolled at different temperatures. After being isothermally held for 2 h, the specimens were cooled to 500 °C at a rate of 30 °C/h, and then cooled in air. Microstructural observations were made optically, with an S-4200 field emission scanning electron microscope (SEM) and with an H-800 transmission electron microscope (TEM). Thin foils for transmission electron microscopic studies were prepared from 3 mm disks, grounded to a thickness of about 50 μm and electropolished in an electrolyte containing 5% perchloric and 95% ethanol at −20 °C. Bright field and dark
field imaging and selected area diffraction patterns were used to identify the carbides. The specimens for Vicker's hardness testing were performed on a Vickers hardness tester (Model: FR-3E) with 30 kg load and each measurement was tested for 8 s. The average hardness of five measurements was obtained as the final hardness.
3.
Results and Discussion
3.1. Microstructure of the AS-sprayed V4 Steel and Phase Transformation Temperature An optical micrograph of the as-sprayed microstructure is shown in Fig. 1a. Spray forming resulted in a substantial reduction in microstructural scale for both grain size and carbide size compared with conventionally cast equivalent [11]. Fine, uniformly distributed equiaxed grains in diameters ranging from 8 to 10 μm are observed. In addition, large primary and interconnected eutectic carbides observed in the conventional cast material are replaced by a more uniformly distributed spheroidal carbides in the as-sprayed material. As shown in Fig. 1b, observation by TEM shows that the matrix of the microstructure of the as-sprayed material was predominantly twinned martensite. Selected area diffraction patterns confirmed the presence of V-rich MC and Cr-base M7C3 type
Fig. 1 – (a) Optical micrograph of the as-sprayed V4 steel, (b) TEM morphology of the twined matrensite.
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teresting to find that an obvious flexure appears during heating between temperatures ranging from 650 °C to 760 °C, indicating the decomposition of martensite. The cooling curve indicates that the high alloyed austenite has a strong stability, as no obvious transformations could be observed other than the bainite reaction which occurred only at very slow cooling rates.
3.2.
Fig. 2 – Measured dilatation variation with temperature.
carbides. Most of the MC and some big M7C3 carbides ranging from 0.5 to 2 μm in diameter are distributed on the grain boundaries. Other small M7C3 carbides with the diameter of about 180 nm distribute in the grains. The variation in specimen dimensions with temperature during heating and cooling process by the dilatometry experiment is shown in Fig. 2. The transformation temperatures, 845 °C (Ac1s) and 890 °C (Ac1f), are defined as the points where the curve leaves the tangent to the linear portion. It is in-
Microstructure Evolution During Hot Rolling
Fig. 3 shows the SEM micrographs of the microstructures before hot rolling at different temperatures. It can be seen that precipitation occurred in all of the cases and obvious differences lie in the amount and morphology of the precipitates. Fig. 3 indicates that the amount of the small carbides decreases with increasing temperature. As shown in Fig. 3a, a larger number of small carbides appear at 850 °C. TEM observation shows that they are M3C or M7C3, however, the bigger, irregular phases, as shown by arrows, are MC (VC) carbides. When the temperature is elevated to 950 °C or even higher, only MC and M7C3 can be found. Thus, reheating temperature leads to major microstructural changes involving phase transformation, dissolution and spheroidization of carbides formed in the as-sprayed steels. Fig. 4 shows the SEM observation on the microstructures of the as-rolled steels which were rolled at 850 °C, 950 °C, 1050 °C, and 1150 °C, respectively. The stable MC carbide particles
Fig. 3 – SEM micrographs of the as-sprayed steels quenched at: (a) 850 °C; (b) 950 °C; (c) 1050 °C and (d) 1150 °C.
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Fig. 4 – SEM micrographs of the as-sprayed steel rolled at: (a) 850 °C; (b) 950 °C; (c) 1050 °C and (d) 1150 °C.
show little tendency to coarsen. The M7C3 carbide, unlike MC, is very sensitive to the rolling temperature, i.e., both the size and morphology of the M7C3 carbides vary with temperature. For example, TEM observation suggests that many coarse carbides shown in Fig. 4a are of M7C3 type. But their size and amount are all decreased with increased rolling temperature. Fig. 5a is the TEM image of those small carbides shown in Fig. 4a. It was confirmed by electron diffraction patterns that most of the small carbides are of M7C3 type, and some M3C carbides are also found. Electron diffraction patterns of M3C and M7C3 carbides are shown in Fig. 5b and c, respectively. The characteristics shown in Fig. 3a and Fig. 4a suggest that the microstructural evolution process at 850 °C differs significantly from those at other temperatures. This can be attributed to the difference of the diffusion coefficient of the alloy elements. Nucleation accompanied by growth and followed by coarsening constitutes the processes during the hot rolling and subsequent cooling process. The nucleation rate as the derivative of the precipitate density N vs time t can be expressed as [12]: s dN DG⁎ 4kR⁎2 DXC0 ¼ N0 Zb⁎ exp and b⁎ ¼ exp : dt t kt a4
ð1Þ
Where N0 is the number of nucleation sites per unit volume, Z is the Zeldovich factor, τ is the incubation time, ΔG⁎ is the nucleation driving force, R⁎ is the nucleation radius,
D is the diffusion coefficient of solute atoms in the matrix, and XC0 is the initial solute mole fraction. The coarsening of second phase particles as a result of a reduction in interfacial energy was theoretically treated by Lifshitz and Slyozov [13] and Wagner where the result was expressed as: ¯r 3 ¯r 30 ¼
2 8gDce Vm t: 9RT
ð2Þ
Where r̄ is the average particle radius at time t, r̄0 is the average particle radius at the onset of coarsening, γ is the interfacial free-energy of the particle–matrix interface, D is the diffusion coefficient of solute atoms in the matrix, ce is the concentration of solute in the matrix in equilibrium with a particle of infinite size, Vm is the molar volume of the particle, R is the gas constant, and T is the temperature. The nucleation rate and coarsening of the carbides are all governed by the diffusion coefficient D of solute atoms, and therefore, the magnitude of the diffusion coefficients provides a useful comparison basis. The diffusion coefficient D can be presented in terms of an Arrhenius equation: Q : D ¼ D0 exp RT
ð3Þ
Where D0 is the pre-experimental factor for the diffusion coefficient, and Q is the activation energy for diffusion.
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to the tempering of martensite around 700 °C [15]. In the early stages of the decomposition, the reaction is controlled by carbon diffusion and is thus very rapid. The change of the material's volume occurs because a large number of M3C carbides precipitate from matrix. This is why an obvious flexure appears in dilatability–temperature curve (Fig. 2) between the temperatures of 650 °C and 760 °C. The amount of the precipitates below the A1 temperature was tremendous, as shown in Fig. 3a. A metastable supersaturated phase will not only generate nuclei but also cause deposition on the generated nuclei and consequent growth of the clusters. This decreases the supersaturation, and correspondingly increases the critical nucleus size, and therefore promotes the coarsening process [16]. As the result of this, the Gibbs–Thomson effect becomes so significant that smaller clusters dissolve, transfer their mass to larger growing clusters and coarsen the size distribution [17]. When small clusters shrink to their critical nucleus size, they become thermodynamically unstable, and spontaneously disintegrate and eventually vanish [18]. This is why a large number of clusters nucleate and then grow, Fig. 3a, and many coarse carbides appear, Fig. 4a. Occasionally coarse carbides, which are favorably situated at the grain corners, appear to have grown from those small carbides close to them. This conclusion is supported by the observation that such particles are frequently polygonal in shape, and there are comparatively fewer small carbides that
Fig. 5 – (a) TEM micrograph of the carbides in the material rolled at 850 °C, electron diffraction patterns in [100] zone axis of M3C particle (b), and in [431] zone of M7C3 particle (c).
According to Eq. (3), the diffusion coefficients of carbon (with D0 = 0.738 cm2 · s− 1 and Q = 158.98 kJ · mol− 1 [14]) and chromium (with D0 = 4.08 cm2 · s− 1 and Q = 286.8 kJ · mol− 1 [14]) at 850 °C are 2.97 × 10− 8 cm2 · s− 1 and 1.86 × 10− 13cm2 · s− 1, respectively. The diffusion coefficients between carbon and chromium are of great difference. Since martensite is not an equilibrium phase, when steel is heated below the eutectoid temperature, the thermodynamically stable phases α and M3C start to precipitate. At temperatures ranging from 650 °C to 760 °C, precipitation of M3C carbides can be written as: α′ ⇒ α + M3C. This reaction is similar
Fig. 6 – (a) TEM micrograph of the polygonal carbide, (b) its electron diffraction pattern in [312] zone axis.
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exist in the vicinity of the coarse particles. On the other hand, in many cases, the firstly formed carbides are non-equilibrium phase and they are replaced by more stable forms gradually in the precipitation sequence [19]. After the precipitation of M3C carbides, the kinetics of the diffusion mechanism changes from carbon diffusion control to chromium diffusion control. Therefore, the observed small M7C3 carbides, as shown in Fig. 5, may be attributed to the transformation from M3C to M7C3 carbides. The possibility of M7C3 carbide formation from M3C carbides was also proved by Homolová [20] and Liu [21,22]. During the heat up process, the carbide precipitation and dissolution is a competition process. At relatively lower temperature the precipitation dominates the process while at high temperature dissolution dominates. Therefore, the decrease of the amount of the small carbides can be attributed to the dissolution with the increase of temperature, and the saturation of the alloy elements in the matrix will increase also. The microstructure of the steel which rolled at 1050 °C shows the best uniform distribution of the carbides with ideal average size, Fig. 4c. Some big polygonal carbides appear at the triplejunctions of grain boundaries in the material which was rolled at 1150 °C, Fig. 4d. Those polygonal carbides were not found in the materials rolled at other temperatures. TEM morphology of this kind of carbide and its corresponding electron diffraction pattern indicate that this carbide is of M7C3, Fig 6. After hot rolling, the supersaturated alloying elements will precipitate during cooling. The precipitation of the carbides is determined by two basic factors. One is that the alloying elements in the matrix should be in supersaturation state, and the other that the atoms have enough ability to diffuse. Both of these factors closely correlate with the temperature. It is thus reasonable to infer that the supersaturation and the diffusion abilities of the alloy atoms are at their highest at 1150 °C in comparison with those at other temperatures. In many diffusion phase transformation and precipitation process, nucleation of the product phase occurs heterogeneously at some preferential nucleation sites in the matrix such as the grain boundaries, dislocations and second phases. Grain boundaries often contain a high density of defects, which can be increased by plastic deformation [23]. Therefore,
Fig. 7 – Hardness of the as-sprayed steels rolled at different temperatures.
Fig. 8 – TEM micrograph of martensite in the as-rolled steel.
carbide precipitation on grain boundaries will be more easily, and precipitates which are growing on boundaries will gain a bigger proportion of their material from within the grains. Although the dissolution of chromium carbides in matrix has been reported by researchers as a more rapid process than the dissolution of other carbides, the complete dissolution of M7C3 carbides is very difficult even at very high austenitizing temperature [24]. When the as-sprayed material was heated to 1150 °C, many M7C3 carbides on the grain boundary still not dissolved completely. They act as preferential nucleation sites and grow up quickly. Thus, big polygonal M7C3 carbides formed and the various shapes of grain boundary M7C3 carbides may possibly be explained by variation of high angle boundaries. There is evidence from Fig. 4d that the most frequently occurring form of grain boundary precipitate is discontinuous. The limited boundary movement observed can be explained in the manner that the precipitation of the carbides in the grain boundaries act as effective obstacles of the growth of the grains, which ensures fine grains remained in elevated temperature. Results of the hardness as a function of the rolling temperature are shown in Fig. 7. It can be seen that the hardness increases with increased rolling temperature. Especially there is a sharp increase from 900 °C to 950 °C. The origination of this effect can be ascribed to the completely different matrix microstructure after hot rolling. TEM observation shows the matrix of the as rolled steels rolled at 850 °C and 900 °C consists of ferrite, as shown in Fig. 5a. Whereas the matrix of the as rolled steels which rolled at or higher than 950 °C consists of martensite, and the representative morphology of the matrix is shown in Fig. 8 (by the steel rolled at 1050 °C). The as-rolled microstructure was controlled by the phase transformation during hot rolling and cooling process, and the rolling temperature is actually a key in this process. When the steel is heated to 900 °C (slightly above A1), little carbide dissolves and the contents of both carbon and alloying elements in austenite are lower. The unstable austenite could transform into ferrite during cooling. But when the temperature is elevated to 950 °C or higher, more alloy elements dissolve in the austenite and improve the hardenability of the steel. Martensite can be therefore form even
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when the steel is cooled in sand. As a result, the completely different microstructure, ferrite and martensite, cause the sharp increase of the hardness from 900 °C to 950 °C.
3.3.
Microstructure Evolution During Annealing
Table 1 – Hardness of the as-rolled steels annealed at different temperatures (HV) Rolling temperature (°C)
Annealing temperature 850 °C
900 °C
950 °C
288 268 278 289 280
243 240 237 246 240
231 230 228 238 219
850 °C
All the as-rolled steels were annealed at different temperatures in order to find the optimum combination of hot rolling and annealing parameters. The best microstructure, with ideal average size and uniform distribution of the carbides, is shown in Fig. 9a. It was obtained by the as-sprayed steel rolled at 1050 °C and then annealed at 900 °C, whereas all other asannealed steels showed inhomogeneous distribution of carbide size to some extent. It can be seen from Fig. 9 that the average carbide size in the spray formed steel is even smaller than the equivalent in the commercial PM V4 steel. TEM observation shows that the as-annealed steel consists of ferrite and homogeneously distributed granular MC and M7C3 carbides. Table 1 compares Vickers hardness of all the asannealed steels. The hardness of those steels which annealed at 900 °C is close to that of the commercial PM V4 steel (HV: 243). The results give strong evidence that the high alloyed V4 steel can be produced by spray forming. Annealing of the as-rolled steels modifies the microstructure in three ways: (1) recrystallization of the matrix; (2) transform the matrix into soft ferrite; and (3) coarsening and further spheroidization of carbides. Microstructural
Fig. 9 – Comparison of V4 annealed microstructures: (a) spray forming, (b) powder metallurgy.
900 °C 950 °C 1050 °C 1150 °C
evolution processes during annealing varied according to the variations of the as rolled steels' matrix. As to those materials, which were rolled at or higher than 950 °C, their matrix is composed of martensite. Similar to the reheating process of the as-sprayed material, when they were heated to elevated temperature the martensite will decompose into ferrite and M3C carbides quickly. But this does not happen during the reheating process to the steels rolled at 850 °C and 900 °C as their matrix is ferrite. In chromium steel the M3C carbides usually have stronger tendency to dissolve in matrix than M7C3 carbides because the later is thermodynamically more stable [24]. Upon heating up to austenitizing temperature, there is a likelihood of in-situ transformation from undissolved M3C carbides to M7C3 carbides. The transformation may begin with the dissolution of M3C followed by in situ nucleation and growth of M7C3 carbides [24], or may proceed by the eutectoid reaction of M3C ⇒ γ + M7C3 [22]. It is well known that a distribution of immobile particles in a solid matrix tends to lower its interfacial free energy by transport of matter from the small to the larger members, thereby diminishing the total particles number but increasing the average particle size. This coarsening process also is known as Ostwald ripening [25,26]. As to the MC and M7C3 carbides, V-rich MC carbides have little tendency to coarsen but Cr-base M7C3 carbides are prone to grow up at elevated temperature. In addition, the growth of dispersed particles in a saturated solution by a diffusion mechanism depends on many variables. As has reported by Greenwood [27], the main factors promoting stability (slow growth-rate) are particles of large mean radius, high density, and low molecular weight, with low values of the solute diffusion coefficient. The carbides' size and distribution of the as-annealed steels are influenced by the microstructures of the as-rolled steels. In general, attractive annealed microstructure can be obtained from the as-rolled steels with fine, homogeneously distributed carbides. Two representative as-annealed steels, which were obtained by rolling at 850 °C and 1050 °C respectively and then all annealed at 900 °C, could be taken to explain the difference. As shown in Fig. 4a, the radii of the M7C3 carbides varied greatly. During the isothermal holding process during annealing, the crucial effect of interfacial curvature on energy, and hence on particle size (Gibbs–Thomson effect), causes larger M7C3 carbides to grow at the expense of the smaller particles, so that smaller particles dissolve and eventually vanish. Thus, it is hard to obtain ideal microstructure by the initial
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incongruous size and distribution of the carbides. On the contrary, as to the specimens rolled at 1050 °C, the carbides are highly dispersed and have a uniform size distribution. Then they have higher stability and grow slowly according to Greenwood's conclusion [27]. No polarizing coarsening behavior as discussed above will appear and it is more likely to obtain ideal equal carbides' size and distribution.
4.
Conclusions
1. Spray forming resulted in a significant refinement of both grain size and carbide size of the high alloyed V4 cold work steel. The optimized processing parameters for the assprayed V4 steel are rolling at 1050 °C and then annealing at 900 °C. No interconnected eutectic carbide structure was observed and fine spheroidal MC and M7C3 carbides were uniformly distributed in the matrix. 2. The hot rolling temperature is the key factor in controlling the evolution of type, morphology and distribution of carbides, as well as in controlling the matrix microstructure of the as-rolled steel. The matrix consists of ferrite when the materials rolled at 850 °C and 900 °C. But when the steels were rolled at or higher than 950 °C, martensite will be obtained because of enhanced hardenability by dissolving of alloy elements in austenite. 3. M3C is the only transient carbide precipitate below the A1 temperature during the reheating of the as-sprayed steel. Irregular M7C3 carbides are preferentially precipitated at the grain boundaries during cooling after the steel was rolled at 1150 °C. 4. The size and distribution of carbides in the as-annealed steels are influenced by the microstructures of the as-rolled steels. Attractive annealed microstructure can be obtained from the as-rolled steels with fine, homogeneously distributed carbides.
Acknowledgements This research was financially supported by Shanghai Baosteel Technology Center and Key Laboratory for High Temperature Materials and Tests of Ministry of Education, Shanghai Jiao Tong University. The author would like to thank Dr Wei Wang for many useful suggestions and comments.
REFERENCES [1] Grant PS. Spray forming. Prog Mater Sci 1995;39:497–545. [2] Ojha SN, Jha JN, Singh SN. Microstructural modifications in Al–Si eutectic alloy produced by spray deposition. Scr Metall Mater 1991;25:443–7. [3] Zhang JG, Xu HB, Shi HS, Wu JS, Sun DS. Microstructure and properties of spray formed Cr12MoV steel for rolls. J Mater Process Technol 2001;111:79–84. [4] Rodenburg C, Krzyzanowski M, Beynon JH, Rainforth WM. Hot workability of spray-fromed AISI M3:2 high-speed steel. Mater Sci Eng A 2004;386:420–7.
[5] Spangel S, Matthaei-Schulz E, Schulz A, Vetters H, Mayr P. Influence of carbon and chromium content and preform shape on the microstructure of spray formed steel deposits. Mater Sci Eng A 2002;326:26–39. [6] Mesquita RA, Barbosa CA. Spray forming high speed steel — properties and processing. Mater Sci Eng A 2004;383:87–95. [7] Hanlon DH, Rainforth WM. The rolling sliding wear response of conventionally proceed and spray formed high speed steel at ambient and elevated temperature. Wear 2003;255:956–66. [8] Yang YF, Hannula SP. Sprayformed hot work steels for rapid tooling. J Mater Sci Technol 2003;19:169–72. [9] Upadhyaya A, Mishra NS, Ojha SN. Microstructural control by spray forming and wear characteristics of a Babbit alloy. J Mater Sci 1997;32:3227–35. [10] Dobrzański LA. Effects of chemical composition and processing conditions on the structure and properties of high-speed steels. J Mater Process Technol 1995;48:727–37. [11] Yan F, et al. Microstructure of the spray formed Vanadis 4 steel and its ultrafine structure. Mater Charact 2007, doi:10.1016/j.matchar.2007.04.019. [12] Michel P, Alexis D. Microscopic modeling of simultaneous two-phase precipitation application to carbide precipitation in low-carbon steels. Mater Sci Eng A 2003;360:214–9. [13] Lifshitz IM, Slyozov VV. The kinetics of precipitation from supersaturated solid solutions. J Phys Chen Solids 1961;19:35–50. [14] Yu YL. Metallic theories. Bei Jing: Metallurgical Industry Press; 2003. p. P213–4. [In Chinese]. [15] Shtansky DV, Inden G. Phase transformation in Fe–Mo–C and Fe–W–C steels—I. The structural evolution during tempering at 700 °C. Acta Mater 1997;45:2861–78. [16] McClurg RB, Flagan RC. Critical comparison of droplet models in homogeneous nucleation theory. J Colloid Interface Sci 1998;201:194–9. [17] Madras G, Mccoy BJ. Temperature effects on the transition from nucleation and growth to Ostwald ripening. Chem Eng Sci 2004;59:2753–65. [18] Madras G, Mccoy BJ. Distribution kinetics theory of Ostwald ripening. J Chem Phys 2001;115:6699–706. [19] Kaneko K, Matsumura S, Sadakata A, Fujita K, Moon WJ, Ozaki S, et al. Characterization of carbides at different boundaries of 9Cr-steel. Mater Sci Eng A 2004;374:82–9. [20] Homolová V, Janovec J, Kroupa A. Experimental and thermodynamic studies of phase transformations in Cr–V low alloy steels. Mater Sci Eng A 2002;335:290–7. [21] Liu ZK, Höglund L, Jönsson B, Ågren J. An experimental and theoretical study of cementite dissolution in an Fe–Cr–C alloy. Metall Trans A 1991;22:1745–52. [22] Liu ZK, Ågren J. Morphology of cementite decomposition in an Fe–Cr–C alloy. Metall Trans A 1991;22:1753–9. [23] Adamson JP, Martin JW. The nucleation of M23C6 carbide particles in the grain boundaries of an austenitic stainless steel. Acta Metall 1971;19:1015–8. [24] Wu C, Sahajwalla V, Krauklis P. The effect of austenitizing process on the hardening behaviour of Cr–Mo–Mn–C air-hardening cast tool steel. ISIJ Inter 1996;36:347–53. [25] Speight MV. Growth kinetics of grain-boundary precipitates. Acta Metall 1968;16:133–5. [26] Ardell AJ. On the coarsening of grain boundary precipitates. Acta Metall 1972;20:601–9. [27] Greenwood GW. The growth of dispersed precipitates in solutions. Acta Metall 1956;4:243–8.