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Microstructure and mechanical property of in-situ nano-particle strengthened ferritic steel by novel internal oxidation
Materials Science & Engineering A 609 (2014) 293–299
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical property of in-situ nano-particle strengthened ferritic steel by novel internal oxidation Hao Tang a, Xiaohua Chen b, Mingwen Chen a, Longfei Zuo a, Bin Hou a, Zidong Wang a,n a b
University of Science and Technology of Beijing, Beijing 100083, PR China State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 4 April 2014 Received in revised form 6 May 2014 Accepted 9 May 2014 Available online 16 May 2014
A novel route of fabricating nano-particles strengthened ferrite steel was investigated in this study. Rather than by externally adding nano-oxide powders, we adopted the endogenous method of controlling oxide reaction and solute concentration distribution in the process of deoxidization to obtain a high density of in-situ nano-oxide particles homogeneously dispersed in the ferrite matrix in melt. The microstructure and tensile properties of these materials had been investigated to clarify the interrelation between the composition, microstructure and mechanical properties. Transmission electron microscopy (TEM) analysis indicated that these nano-particles were titanium oxides, which have a positive effect on optimizing inclusions and refining grains. Tensile tests revealed that these titanium oxide particles play an important role in increasing the yield strength. The steel has yield strength of 711 MPa, approximately three times higher than that of conventional plain carbon structural steels, and its ultimate tensile strength reaches 810 MPa with an elongation-to-failure value of 22%. Precipitation hardening and grain refinement hardening are the dominant factors responsible for yield strength increasing in this steel. & 2014 Elsevier B.V. All rights reserved.
Keywords: in situ nano-particles Titanium oxide In melt Precipitation hardening
1. Introduction At present, nano-particles strengthened ferritic steels are typically referred to as oxide dispersion strengthened (ODS) steels by externally adding nano yttrium oxide powders [1]. As summarized by Schneibel et al. [2], the manufacturing of ODS ferritic alloys usually takes a long time (e.g., one day) in a high-energy ball mill, followed by hot consolidation. The mechanical alloying process for creating oxide dispersoid is expensive and energy-intensive. Also, mechanically alloyed materials typically develop pores during high temperature annealing [3]. These pores degrade the mechanical properties. Issues for ODS ferritic or ferritic/martensitic steels processed by mechanical alloying (MA) are anisotropic mechanical properties due to the bamboo-like structure and impurity pick up during MA [4]. Long time high-energy mechanical milling could bring contamination from the atmosphere and the milling media; the effects of such contamination on structure, mechanical properties, and irradiation performance of the MA products could be vital. Furthermore, the nano-particles strengthened steel prepared by adding nano oxide particles externally generates such problems like the uneven distribution or the press that is easily caused by
n
Corresponding author. Tel./fax: þ 86 010 62333979. E-mail address: [email protected] (Z. Wang).
http://dx.doi.org/10.1016/j.msea.2014.05.020 0921-5093/& 2014 Elsevier B.V. All rights reserved.
the interface between the particle and the matrix. Compared to external adding, the endogenous method does not need to prepare nano powder. It is low-cost, time saving and has high efficiency. The in-situ nano-particles prepared in this way are more dispersive, uniform and present a coherent or semi-coherent relationship with the matrix. The formation of coherent or semi-coherent interface will reduce the interfacial energy between oxide particle and matrix, which results in the increasing stability of the oxide particles [5]. The objective of this work is to take advantage of the in-situ nano oxide particles produced by a novel internal oxidation to strengthen the ferritic steel. The manufacturing route, microstructure and mechanical properties of the steel have been preliminarily investigated.
2. Experimental methods Two steels, with and without titanium, were prepared and studied. The ingots were produced by vacuum induction melting furnace and the main chemical compositions of the steel investigated are shown in Table 1. Steel B has the same composition as that of steel A except the addition of 0.15% Ti. The former adopted a special route to produce nano precipitation in melt directly. Here we show the nanoparticles formed in melt directly which is
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H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
Table 1 Main chemical composition of steels investigated (wt%). Steel A
C 0.068
Si 0.13
Mn 0.28
S 0.030
P 0.024
N 0.004
–
O
Steel B
C 0.068
Si 0.13
Mn 0.28
S 0.030
P 0.024
N 0.004
Ti: 0.15
o 40 ppm
Tire wire multipoint dispersion supply
Ti+O
TixOy
electromagnetic stirring convection field Fig. 1. Schematic of the processing method for the formation of nanoparticles in melt through a combination of multipoint dispersion supply processing and electromagnetic stirring.
achieved by solidification processing through a combination of uniform dispersion supply processing and electromagnetic stirring. Fig. 1 schematically shows the specific experimental process. First, the base metal is placed in the vacuum melting furnace. After it fully melts and the solidification temperature of the molten steel is in the range from 1530 1C to 1620 1C, the pure Ti wire is added in the way of multipoint dispersion supply. It is worth noting that electromagnetic stirring ( 4 KHz frequency) always exists during the whole melting process. Under the effect of electromagnetic stirring, flowing linear velocity of the molten metal is about 10–20 m/s. It is followed by thermal insulation for about 60 s after all the pure Ti wires completely react. In the casting process, the flowing linear velocity of the molten metal is not less than 1.7 m/s. Finally, the melt was cooled down inside the crucible in air over a cooling rate of about 500 1C/min in the process of solidification. Subsequently, two ingots with dimensions of 135 mm 75 mm 96 mm were austenitized at 1200 1C for 2 h, then hot rolled into 11 mm thick plate by seven passes. The initial and final rolling temperatures were 1150 1C and 870 1C, respectively. Finally the specimens were water-cooled to a certain temperature ( 600 1C) and then air-cooled to ambient temperature. Metallographic specimens were cut perpendicular to the roll direction from hot plate and were ground, polished, etched with 3% nitric acid–alcohol and then observed with 9XB-PC optical microscopy. Microstructure of the specimen was further investigated by using JEOL JEM-2100 TEM operated at 120 kV. TEM foil was prepared by cutting a thin wafer from the steel samples, followed by mechanically thinning them to 35 μm in thickness. Three millimeter discs were punched from the wafers and electrochemically-polished using a solution of 10 vol% HClO4– methanol electrolyte at low temperature. The foils were thinned further for electron transparent area using focused ion beam. According to ASTM E8 specification, dog-bone shaped plate on longitudinal specimens with a gauge section of 30 mm Length 20 mm Width 2 mm Thickness were used for the tensile tests,
which were conducted at room temperature with a strain rate of 10 3 s 1 using an FPZ 100 machine.
3. Results 3.1. Microstructure Fig. 2a shows the morphology and distribution of the inclusions using TEM. Many cubic inclusions about several hundred nanometers can be observed in steel A. However, as Fig. 2b shows, a remarkable decrease in the inclusion number can be observed in steel B. The inclusion shape changes from cubic into circular and it locates in the intersections of grain boundary. It is easy to see the inclusion pinning on the grain boundary which will stop the grains from growing up. Energy Disperse Spectroscopy (EDS) is used to determine the constitution of the inclusion in the two steels. Results indicate that the cubic inclusions are MnS while the circular one is a composite inclusion of MnS and Titanium Oxide. Fig. 3 shows the optical micrographs of the investigated steels. It can be seen that the average grain size of steel B is obviously smaller than that of steel A, due to the comprehensive function of the nano precipitated phase. 3.2. Characterization of nano precipitates Fig. 4 shows representative TEM micrographs in the investigated steel. Without adopting special craft, there are no nanoparticles in steel A, as shown in Fig. 4a. On the contrary, it is obvious to see that a large number of nano-particles are homogeneously distributed in the matrix with mean diameter of approximately 5 nm and sizes smaller than 10 nm in steel B, as shown in Fig. 4(b, c). Fig. 4d shows the corresponding darkfield image.
H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
295
200 nm
100 nm Fig. 2. TEM micrograph and EDS result of inclusions in steels investigated: (a) steel A and (b) steel B.
60 µm
60 µm Fig. 3. Low magnification light micrographs of steels investigated: (a) steel A and (b) steel B.
The HRTEM image of Ti–O nano-particles (bcc-Fe [001] zone axis) in the alloy is shown in Fig. 5. Inset shows the FFT pattern (the additional reflections are circled) from the corresponding particle; the white rectangle shows the reflection with a d-spacing
of 0.3037 nm, corresponding to the (301) plane (d spacing 0.3048 nm) of Ti3O5. Once the nano titanium oxide is formed in the melt, they have a strong tendency to grow upwards. So this process is controlled
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H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
100 nm
100 nm
5 nm
50 nm
50 nm
Fig. 4. TEM images of the steels investigated: (a) steel A. (b, c) A bright-field image of steel B and (d) the corresponding dark-field image.
strictly. Fig. 6 shows the influence of factors for the growth speed of nano-particle. The growth speed of the nano particle is concerned with the [Ti] content in the front of the interface and its radius. Obviously we can see that with the [Ti] content reducing from 100% to 0.01%, the growth speed also reduces sharply. At the level of 0.01%, its growth speed is almost zero. Also notice that with the increase of the nano-particle's radius, its growth speed is slower. So, the growth speed is influenced by the [Ti] content in the micro area and the radius of the nano-particle. By controlling these parameters precisely a large number of nanoscale particles can be obtained in the melt finally.
precipitates tangled with dislocations can be found in Fig. 7c. The dislocation lines, pinned by precipitate particle, bow out to surmount plastic deformation, as indicated by the arrows. 3.4. Mechanical properties Fig. 8 shows the corresponding engineering stress–strain curve for steel B. As shown in Table 2, compared with steel A, tensile strength and yield strength of steel B are increased from 372 MPa and 272 MPa to 810 MPa and 711 MPa, by 117% and 167%, respectively. While the plasticity and toughness do not decrease too much.
3.3. Dislocations under rolling condition 4. Discussion Fig. 7 shows dislocations of the configuration under the rolling condition. We can see in Fig. 7a that there are only a few dislocations and no nano precipitates in steel A. However, in Fig. 7b, a remarkable increase in the dislocation density can be observed in steel B and it shows the interaction between nano precipitates and dislocations in the matrix. Finely dispersed
4.1. Inclusions optimization It is well established that different internal factors of inclusion have a great influence on the performance of steel. Various defects can be generated, such as the crack initiation site under external
H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
d=0.3037nm
200 2
1 110
3 103
110 6 4
5
211
114 123
αFe
456
Ti3O5
5 nm Fig. 5. The HRTEM image of Ti–O nano-particles (along the bcc-Fe [001] direction) in steel B. Inset shows the FFT pattern (the additional reflections are circled) from the corresponding particle; the white rectangle shows the reflection with a dspacing of 0.3037 nm, which corresponds to the (301) plane (d spacing 0.3048) of the Ti3O5.
kinds of measurements have been taken to remove these inclusions. However, the above mentioned “Oxides Metallurgy Technology” changes the traditional viewpoint, if we can control the internal factors of inclusion such as reducing its size and controlling its distribution in steel, etc. This will ameliorate the influence of the inclusion to a maximum extent. Compared to sulfur, the properties of oxide in the steel are easier to control. In our experiment, a large number of dispersive nano oxides were formed in the melt. These preferentially formed nano-particles will be the heterogeneous nucleus of the sulfide; thus the distribution, size and the shape of the sulfide are under good control [10,11]. The number and size of inclusions significantly decreased and their shape significantly optimized. As shown in Fig. 2b, the size of the composite inclusion is less than one hundred nanometers and it is distributed in the intersections of the grain boundary. Obviously, the optimized inclusion pinning at the grain boundary will restrain the growth of grains and generate grain refinement hardening. 4.2. Formation of the nano titanium oxide in melt Titanium exhibits a strong tendency to form oxides and titanium compounds will precipitate from the steel during the casting process. When adding titanium wire, its tip starts to melt, generates a large number of free titanium [Ti] and quickly diffuses to the surrounding melt. Micro area concentration gradient of [Ti] will be set up around the tip of the bulk pure Ti in a short time. Meanwhile, Ti will react with oxygen and consume the free-state oxygen [O] in the tip area. A large number of titanium oxide homogeneous nucleation cores are formed in this process. Nano oxide nuclear that grow upwards belongs to the typical solute element diffusion growth, whose growth speed is influenced by the size of the oxide and the solute micro area concentration fluctuations in the front of the oxide interface. Its growth diffusion equation can be predicted by the following equation [12]: dr M s ρm ¼ DL ðC L C e Þ dt 100M m r ρs
Fig. 6. Radius of the nano titanium oxide particle vs. growth speed under the condition of different [Ti] concentrations at the temperature of 1873 K.
load, if there are lager-size brittle inclusions [6]. MnS is one of the most common nonmetal plastic inclusions in steel. Its size, shape and distribution seriously affect the performance of steel [7–9]. All
297
ð1Þ
where r is the diameter of the oxide; DL is the solute diffusion coefficient in liquid steel; ρ is the density; Ms is the molecular weight of oxide; Mm is the molecular weight of alloy element; CL is the solute concentration in the front of the oxide interface; and Ce is the equilibrium solute concentration of the oxide. It is obvious to see that the growth driving force of the nano particle is the solute segregation of the solute atom in the front of the oxide interface and the equilibrium concentration of the oxide (CL(t)–Ce(t)). Its growth speed is influenced by the size of the oxide and the solute micro area concentration fluctuations in the front of the oxide interface. As Fig. 6 shows, the growth speed of the nano particle is slower with the increase of its radius; the growth speed of the nano particle reduced sharply, even reduced to zero, with the decrease of the [Ti] content in the front of the interface. We can conclude from Fig. 6 and Eq. (1) that growth speed of the nano particle is proportional to the difference value (CL(t)–Ce(t)), namely, under a certain temperature, it is only influenced by the [Ti] content in the micro area. We can control its growth speed in the liquid steel by adjusting the two parameters. So, as the titanium oxide homogeneous nucleation core is formed in melt, simultaneously, the electromagnetic stirring is carried out to generate the strong convection field in the metal melt so that the new formed titanium oxide homogeneous nucleation core will be taken away as soon as possible. Definitely, the convection coefficient must be strictly controlled. The strong convection field will break away the micro area concentration gradient of [Ti] in order to lower or cut off the [Ti] concentration that surrounds the nano nucleation core, namely, control the value of CL. So, if there is
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H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
200 nm
200 nm
100 nm Fig. 7. TEM micrograph showing dislocation configuration under rolling condition in steels investigated: (a) steel A and (b) steel B. (c) TEM image of steel B shows the interaction between nano oxide precipitation and dislocation, which is indicated by arrows.
Table 2 Longitudinal mechanical properties of steels investigated.
Steel A Steel B
Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
272 711
372 810
33 22
melt directly. Relevant information on the materials fabrication can be found elsewhere [13]. 4.3. Strengthening mechanism The relationship between yield strength and microstructure can be predicted by the following equation [14]:
ss ¼ s0 þ sg þ sp þ sss þ sd Fig. 8. Tensile stress–strain curves for steel B.
no concentration supply, the further growth of the nano particle will not meet thermodynamics and dynamics conditions and their size will always remain in a smaller size range. Through this process, we can obtain a large number of nanoscale particles in
ð2Þ
where s0 is the internal lattice strength; sg , sp , sss and sd are the strengthening effects caused by grain refinement, precipitation, solid solution, and dislocation, respectively. For present interest, the hardening from precipitates and grain refinement plays a dominant role in steel B. The other strengthening mechanism cannot be overlooked either.
H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
More recently, it was observed that nanoparticles can effectively serve as a catalyst for heterogeneous nucleation of metal grains during solidification [15,16] and restrictions of grain growth by the effect of Zener pinning [17–19] during solidification and solid-state heat treatment for the existence of a large number of nanoparticles distributed in the grain boundary and intracrystalline. In this work, the nanoparticles formed in melt can also be the heterogeneous nucleus of the sulfide, thus to optimize the inclusion. Obviously, as shown in Fig. 2b, the optimized inclusion pinning at the grain boundary will restrain the growth of grains and generate grain refinement hardening. All the above comprehensive functions of the nano precipitated phase result in significantly smaller grain sizes and markedly improved yield strength according to the Hall–Petch relationship. As shown in Fig. 3, the average grain size of steel B is obviously smaller than that of steel A. As described by the Hall–Petch equation [20]:
s ¼ K y d 1=2
ð3Þ
where d is the average ferrite grain size; Ky is a constant, 15:1–18:1 N mm 3=2 . The average ferrite grain size has been estimated using the Image-Pro Plus software. For the steel with an average grain size of 6 μm, a grain refinement hardening effect of about 225 MPa can be expected. The dominant factor for precipitation hardening is generated by a large number of dispersing nanometer particles for this new steel B when compared with steel A. Precipitation hardening lies in the interaction between the second precipitation and the dislocations, causing the blocking of the dislocation movement, contributing to the high tensile strength due to the Orowan hardening mechanism [21,22]. In Fig. 7a, b, a significant change in the microstructure of the steel occurs. Compared with steel A, a remarkable increase in the dislocation density can be observed, suggesting that the finely dispersed precipitates are beneficial for the increase in dislocation density. Furthermore, it is obvious to see that most of the precipitates are tangled with dislocations, as shown in Fig. 7(c). The HRTEM of precipitated Ti3O5 nano-particles indicates that the lattice contiguity between the oxide particles and matrix is evident. While we substantially improve its strength, the coherent or semi-coherent relationship guarantees that when the moving dislocations meet these coherent particles, part of them will pile up and others will continue to keep slipping along the coherent interface [23]. And this will not lead to the aggregation and hinder of dislocations on the grain boundaries; thus we can improve the strength without sacrificing its ductility too much. According to Gladman's theory, using the modified Ashby– Orowan model to calculate the precipitation hardening effect [24,25]: Rp ¼
10μb r ϕ1=2 ln b 5:72r π 3=2
ð4Þ
where μ is the shearing factor (as for ferrite, μ ¼ 80.26 103 MPa); b is the Burgers Vector, 2.48 10–4 μm; r is the radius of the particle;ϕ is the volume fraction of the precipitates. The precipitate radius (r) and the volume fraction of the precipitates (ϕ) have been estimated using small-angle X-ray-scattering (SAXS) and Fit2d software, which are about 6.7 nm and 0.354%, respectively. Consequently, precipitation hardening generated by nano precipitates is determined using Eq. (4) as 183 MPa. Once again, the above value reveals that the finely dispersed precipitates which are formed in melt can not only have a positive effect on precipitation hardening, but also can act as heterogeneous nucleation core thus to refine as-cast structure grains and optimize inclusions.
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5. Conclusions A novel route of fabricating in-situ nano-particles strengthened ferritic steel was investigated in this study. The main conclusions can be drawn as follows: (1) A ferritic steel reinforced with in-situ nanoparticles which have mean diameter of approximately 5 nm and sizes smaller than 10 nm has been fabricated by a novel route. (2) Size of the inclusions which are nucleated on the preferentially formed nano-particles is refined greatly and their distribution is under good control. The inclusion pinning on the grain boundary will restrain the growth of grains. (3) The nano-particles in the investigated steel are Ti3O5. These particles share a coherent or semi-coherent relationship with the matrix. Thus we can improve its strength substantially without sacrificing its plasticity and toughness too much. (4) The contribution of this novel fabricating route to yield strength reaches 439 MPa in the high strength steel. Precipitation of nano-particles is the dominant factor responsible for yield strength increasing.
Acknowledgments The authors are greatly indebted to Prof. Zidong Wang and Dr. Xiaohua Chen for valuable discussions. We also thank Dr Longfei Zuo for his participation of part of the experiments. This work was supported by the Specialized Research Fund for the Doctoral Program of Higher Education of China (No. 20100006110026). References [1] K. Ehrlich, Fusion Eng. Des. 56 (2001) 71–82. [2] J.H. Schneibel, S. Shim, Mater. Sci. Eng. A 488 (2008) 134–138. [3] J.H. Schneibel, C.T. Liu, D.T. Hoelzer, M.J. Mills, P. Sarosi, T. Hayashi, U. Wendt, H. Heyse, Scr. Mater. 57 (2007) 1040–1043. [4] S. Yamashita, S. Watanabe, S. Ohnuki, H. Takahashi, N. Akasaka, S. Ukai, J. Nucl. Mater. 283 (2000) 647–651. [5] L.S. Zhang, S. Ukai, T. Hoshino, S. Hayashi, X.H. Qu, Acta Mater. 57 (2009) 3671–3682. [6] Z.Z. Liu, Y.S. Sang, Steelmaking 23 (2007) 3. [7] W. Yang, X.G. Yang, L.F. Zhang, F.X. Liu, Steelmaking 29 (2013) 6. [8] Y.M. Li, F.X. Zhu, F.P. Cui, J. Northeast. Univ. 28 (2007) 1002–1005. [9] G. Domizzi, G. Anteri, J. Ovejero-Garcia, Corros. Sci. 43 (2001) 325–339. [10] J.I. Takamura, S. Mizoguchi, Metallurgy of oxides in steels. I. Roles of oxides in steels performance[C]//IISC, in: Proceedings of The Sixth International Iron and Steel Congress vol. 1 (1990), pp. 591–597. [11] T. Sawai, M. Wakoh, S. Mizoguchi, J. Iron Steel Inst. Jpn. 82 (7) (1996) 587–592. [12] H. Goto, K. Miyazawa, K. Amaguchi, et al., ISIJ Int. 34 (1994) 414–419. [13] Z.D. Wang, X.W. Wang, Q.S. Wang, Nanotechnology 20 (2009) 075605. [14] F.B. Pickering, Physical metallurgy and the design of steels[M], Applied Science Publication Limited, London, 1978, p. 63. [15] M.P. De Cicco, L.S. Turng, X.C. Li, J.H. Perepezko, Metall. Mater. Trans. A 42 (2011) 2323–2330. [16] A. Belyakov, F.G. Wei, K. Tsuzaki, Y. Mishima, Mater. Sci. Eng. A 471 (2007) 50. [17] E. Nes, N. Ryum, O. Hunderi, Acta Mater. 33 (1985) 11–22. [18] E. Nes, N. Ryum, O. Hunderi, Acta Metall. 33 (1985) 1. [19] U.F. Kocks, H. Meching, Acta Metall. 29 (1981) 1865. [20] X.P. Mao, X.D. Huo, X.J. Sun, Y.Z. Chai, J. Mater. Process. Technol. 210 (2010) 1660–1666. [21] Y.R. Wen, A. Hirata, Z.W. Zhang, T. Fujita, C.T. Liu, J.H. Jiang, M.W. Chen, Acta Mater. 61 (2013) 2133–2147. [22] W.D. Callister, Materials Science and Engineering: An Introduction, Wiley, New York, 2007. [23] C.R. Dong, H.P. Reng, T.Z. Jing, Microalloyed Non Quenched and Tempered Steel, Metallurgical Industry Press, Beijing, 2000. [24] Y.L. Kang, H. Yu, Steel 38 (2003) 20–26. [25] T. Gladman, Miner. Met. Mater. Soc. 3 (1992) 14.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Microstructure and mechanical property of in-situ nano-particle strengthened ferritic steel by novel internal oxidation Hao Tang a, Xiaohua Chen b, Mingwen Chen a, Longfei Zuo a, Bin Hou a, Zidong Wang a,n a b
University of Science and Technology of Beijing, Beijing 100083, PR China State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 4 April 2014 Received in revised form 6 May 2014 Accepted 9 May 2014 Available online 16 May 2014
A novel route of fabricating nano-particles strengthened ferrite steel was investigated in this study. Rather than by externally adding nano-oxide powders, we adopted the endogenous method of controlling oxide reaction and solute concentration distribution in the process of deoxidization to obtain a high density of in-situ nano-oxide particles homogeneously dispersed in the ferrite matrix in melt. The microstructure and tensile properties of these materials had been investigated to clarify the interrelation between the composition, microstructure and mechanical properties. Transmission electron microscopy (TEM) analysis indicated that these nano-particles were titanium oxides, which have a positive effect on optimizing inclusions and refining grains. Tensile tests revealed that these titanium oxide particles play an important role in increasing the yield strength. The steel has yield strength of 711 MPa, approximately three times higher than that of conventional plain carbon structural steels, and its ultimate tensile strength reaches 810 MPa with an elongation-to-failure value of 22%. Precipitation hardening and grain refinement hardening are the dominant factors responsible for yield strength increasing in this steel. & 2014 Elsevier B.V. All rights reserved.
Keywords: in situ nano-particles Titanium oxide In melt Precipitation hardening
1. Introduction At present, nano-particles strengthened ferritic steels are typically referred to as oxide dispersion strengthened (ODS) steels by externally adding nano yttrium oxide powders [1]. As summarized by Schneibel et al. [2], the manufacturing of ODS ferritic alloys usually takes a long time (e.g., one day) in a high-energy ball mill, followed by hot consolidation. The mechanical alloying process for creating oxide dispersoid is expensive and energy-intensive. Also, mechanically alloyed materials typically develop pores during high temperature annealing [3]. These pores degrade the mechanical properties. Issues for ODS ferritic or ferritic/martensitic steels processed by mechanical alloying (MA) are anisotropic mechanical properties due to the bamboo-like structure and impurity pick up during MA [4]. Long time high-energy mechanical milling could bring contamination from the atmosphere and the milling media; the effects of such contamination on structure, mechanical properties, and irradiation performance of the MA products could be vital. Furthermore, the nano-particles strengthened steel prepared by adding nano oxide particles externally generates such problems like the uneven distribution or the press that is easily caused by
n
Corresponding author. Tel./fax: þ 86 010 62333979. E-mail address: [email protected] (Z. Wang).
http://dx.doi.org/10.1016/j.msea.2014.05.020 0921-5093/& 2014 Elsevier B.V. All rights reserved.
the interface between the particle and the matrix. Compared to external adding, the endogenous method does not need to prepare nano powder. It is low-cost, time saving and has high efficiency. The in-situ nano-particles prepared in this way are more dispersive, uniform and present a coherent or semi-coherent relationship with the matrix. The formation of coherent or semi-coherent interface will reduce the interfacial energy between oxide particle and matrix, which results in the increasing stability of the oxide particles [5]. The objective of this work is to take advantage of the in-situ nano oxide particles produced by a novel internal oxidation to strengthen the ferritic steel. The manufacturing route, microstructure and mechanical properties of the steel have been preliminarily investigated.
2. Experimental methods Two steels, with and without titanium, were prepared and studied. The ingots were produced by vacuum induction melting furnace and the main chemical compositions of the steel investigated are shown in Table 1. Steel B has the same composition as that of steel A except the addition of 0.15% Ti. The former adopted a special route to produce nano precipitation in melt directly. Here we show the nanoparticles formed in melt directly which is
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H. Tang et al. / Materials Science & Engineering A 609 (2014) 293–299
Table 1 Main chemical composition of steels investigated (wt%). Steel A
C 0.068
Si 0.13
Mn 0.28
S 0.030
P 0.024
N 0.004
–
O
Steel B
C 0.068
Si 0.13
Mn 0.28
S 0.030
P 0.024
N 0.004
Ti: 0.15
o 40 ppm
Tire wire multipoint dispersion supply
Ti+O
TixOy
electromagnetic stirring convection field Fig. 1. Schematic of the processing method for the formation of nanoparticles in melt through a combination of multipoint dispersion supply processing and electromagnetic stirring.
achieved by solidification processing through a combination of uniform dispersion supply processing and electromagnetic stirring. Fig. 1 schematically shows the specific experimental process. First, the base metal is placed in the vacuum melting furnace. After it fully melts and the solidification temperature of the molten steel is in the range from 1530 1C to 1620 1C, the pure Ti wire is added in the way of multipoint dispersion supply. It is worth noting that electromagnetic stirring ( 4 KHz frequency) always exists during the whole melting process. Under the effect of electromagnetic stirring, flowing linear velocity of the molten metal is about 10–20 m/s. It is followed by thermal insulation for about 60 s after all the pure Ti wires completely react. In the casting process, the flowing linear velocity of the molten metal is not less than 1.7 m/s. Finally, the melt was cooled down inside the crucible in air over a cooling rate of about 500 1C/min in the process of solidification. Subsequently, two ingots with dimensions of 135 mm 75 mm 96 mm were austenitized at 1200 1C for 2 h, then hot rolled into 11 mm thick plate by seven passes. The initial and final rolling temperatures were 1150 1C and 870 1C, respectively. Finally the specimens were water-cooled to a certain temperature ( 600 1C) and then air-cooled to ambient temperature. Metallographic specimens were cut perpendicular to the roll direction from hot plate and were ground, polished, etched with 3% nitric acid–alcohol and then observed with 9XB-PC optical microscopy. Microstructure of the specimen was further investigated by using JEOL JEM-2100 TEM operated at 120 kV. TEM foil was prepared by cutting a thin wafer from the steel samples, followed by mechanically thinning them to 35 μm in thickness. Three millimeter discs were punched from the wafers and electrochemically-polished using a solution of 10 vol% HClO4– methanol electrolyte at low temperature. The foils were thinned further for electron transparent area using focused ion beam. According to ASTM E8 specification, dog-bone shaped plate on longitudinal specimens with a gauge section of 30 mm Length 20 mm Width 2 mm Thickness were used for the tensile tests,
which were conducted at room temperature with a strain rate of 10 3 s 1 using an FPZ 100 machine.
3. Results 3.1. Microstructure Fig. 2a shows the morphology and distribution of the inclusions using TEM. Many cubic inclusions about several hundred nanometers can be observed in steel A. However, as Fig. 2b shows, a remarkable decrease in the inclusion number can be observed in steel B. The inclusion shape changes from cubic into circular and it locates in the intersections of grain boundary. It is easy to see the inclusion pinning on the grain boundary which will stop the grains from growing up. Energy Disperse Spectroscopy (EDS) is used to determine the constitution of the inclusion in the two steels. Results indicate that the cubic inclusions are MnS while the circular one is a composite inclusion of MnS and Titanium Oxide. Fig. 3 shows the optical micrographs of the investigated steels. It can be seen that the average grain size of steel B is obviously smaller than that of steel A, due to the comprehensive function of the nano precipitated phase. 3.2. Characterization of nano precipitates Fig. 4 shows representative TEM micrographs in the investigated steel. Without adopting special craft, there are no nanoparticles in steel A, as shown in Fig. 4a. On the contrary, it is obvious to see that a large number of nano-particles are homogeneously distributed in the matrix with mean diameter of approximately 5 nm and sizes smaller than 10 nm in steel B, as shown in Fig. 4(b, c). Fig. 4d shows the corresponding darkfield image.
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200 nm
100 nm Fig. 2. TEM micrograph and EDS result of inclusions in steels investigated: (a) steel A and (b) steel B.
60 µm
60 µm Fig. 3. Low magnification light micrographs of steels investigated: (a) steel A and (b) steel B.
The HRTEM image of Ti–O nano-particles (bcc-Fe [001] zone axis) in the alloy is shown in Fig. 5. Inset shows the FFT pattern (the additional reflections are circled) from the corresponding particle; the white rectangle shows the reflection with a d-spacing
of 0.3037 nm, corresponding to the (301) plane (d spacing 0.3048 nm) of Ti3O5. Once the nano titanium oxide is formed in the melt, they have a strong tendency to grow upwards. So this process is controlled
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100 nm
100 nm
5 nm
50 nm
50 nm
Fig. 4. TEM images of the steels investigated: (a) steel A. (b, c) A bright-field image of steel B and (d) the corresponding dark-field image.
strictly. Fig. 6 shows the influence of factors for the growth speed of nano-particle. The growth speed of the nano particle is concerned with the [Ti] content in the front of the interface and its radius. Obviously we can see that with the [Ti] content reducing from 100% to 0.01%, the growth speed also reduces sharply. At the level of 0.01%, its growth speed is almost zero. Also notice that with the increase of the nano-particle's radius, its growth speed is slower. So, the growth speed is influenced by the [Ti] content in the micro area and the radius of the nano-particle. By controlling these parameters precisely a large number of nanoscale particles can be obtained in the melt finally.
precipitates tangled with dislocations can be found in Fig. 7c. The dislocation lines, pinned by precipitate particle, bow out to surmount plastic deformation, as indicated by the arrows. 3.4. Mechanical properties Fig. 8 shows the corresponding engineering stress–strain curve for steel B. As shown in Table 2, compared with steel A, tensile strength and yield strength of steel B are increased from 372 MPa and 272 MPa to 810 MPa and 711 MPa, by 117% and 167%, respectively. While the plasticity and toughness do not decrease too much.
3.3. Dislocations under rolling condition 4. Discussion Fig. 7 shows dislocations of the configuration under the rolling condition. We can see in Fig. 7a that there are only a few dislocations and no nano precipitates in steel A. However, in Fig. 7b, a remarkable increase in the dislocation density can be observed in steel B and it shows the interaction between nano precipitates and dislocations in the matrix. Finely dispersed
4.1. Inclusions optimization It is well established that different internal factors of inclusion have a great influence on the performance of steel. Various defects can be generated, such as the crack initiation site under external
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d=0.3037nm
200 2
1 110
3 103
110 6 4
5
211
114 123
αFe
456
Ti3O5
5 nm Fig. 5. The HRTEM image of Ti–O nano-particles (along the bcc-Fe [001] direction) in steel B. Inset shows the FFT pattern (the additional reflections are circled) from the corresponding particle; the white rectangle shows the reflection with a dspacing of 0.3037 nm, which corresponds to the (301) plane (d spacing 0.3048) of the Ti3O5.
kinds of measurements have been taken to remove these inclusions. However, the above mentioned “Oxides Metallurgy Technology” changes the traditional viewpoint, if we can control the internal factors of inclusion such as reducing its size and controlling its distribution in steel, etc. This will ameliorate the influence of the inclusion to a maximum extent. Compared to sulfur, the properties of oxide in the steel are easier to control. In our experiment, a large number of dispersive nano oxides were formed in the melt. These preferentially formed nano-particles will be the heterogeneous nucleus of the sulfide; thus the distribution, size and the shape of the sulfide are under good control [10,11]. The number and size of inclusions significantly decreased and their shape significantly optimized. As shown in Fig. 2b, the size of the composite inclusion is less than one hundred nanometers and it is distributed in the intersections of the grain boundary. Obviously, the optimized inclusion pinning at the grain boundary will restrain the growth of grains and generate grain refinement hardening. 4.2. Formation of the nano titanium oxide in melt Titanium exhibits a strong tendency to form oxides and titanium compounds will precipitate from the steel during the casting process. When adding titanium wire, its tip starts to melt, generates a large number of free titanium [Ti] and quickly diffuses to the surrounding melt. Micro area concentration gradient of [Ti] will be set up around the tip of the bulk pure Ti in a short time. Meanwhile, Ti will react with oxygen and consume the free-state oxygen [O] in the tip area. A large number of titanium oxide homogeneous nucleation cores are formed in this process. Nano oxide nuclear that grow upwards belongs to the typical solute element diffusion growth, whose growth speed is influenced by the size of the oxide and the solute micro area concentration fluctuations in the front of the oxide interface. Its growth diffusion equation can be predicted by the following equation [12]: dr M s ρm ¼ DL ðC L C e Þ dt 100M m r ρs
Fig. 6. Radius of the nano titanium oxide particle vs. growth speed under the condition of different [Ti] concentrations at the temperature of 1873 K.
load, if there are lager-size brittle inclusions [6]. MnS is one of the most common nonmetal plastic inclusions in steel. Its size, shape and distribution seriously affect the performance of steel [7–9]. All
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ð1Þ
where r is the diameter of the oxide; DL is the solute diffusion coefficient in liquid steel; ρ is the density; Ms is the molecular weight of oxide; Mm is the molecular weight of alloy element; CL is the solute concentration in the front of the oxide interface; and Ce is the equilibrium solute concentration of the oxide. It is obvious to see that the growth driving force of the nano particle is the solute segregation of the solute atom in the front of the oxide interface and the equilibrium concentration of the oxide (CL(t)–Ce(t)). Its growth speed is influenced by the size of the oxide and the solute micro area concentration fluctuations in the front of the oxide interface. As Fig. 6 shows, the growth speed of the nano particle is slower with the increase of its radius; the growth speed of the nano particle reduced sharply, even reduced to zero, with the decrease of the [Ti] content in the front of the interface. We can conclude from Fig. 6 and Eq. (1) that growth speed of the nano particle is proportional to the difference value (CL(t)–Ce(t)), namely, under a certain temperature, it is only influenced by the [Ti] content in the micro area. We can control its growth speed in the liquid steel by adjusting the two parameters. So, as the titanium oxide homogeneous nucleation core is formed in melt, simultaneously, the electromagnetic stirring is carried out to generate the strong convection field in the metal melt so that the new formed titanium oxide homogeneous nucleation core will be taken away as soon as possible. Definitely, the convection coefficient must be strictly controlled. The strong convection field will break away the micro area concentration gradient of [Ti] in order to lower or cut off the [Ti] concentration that surrounds the nano nucleation core, namely, control the value of CL. So, if there is
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200 nm
200 nm
100 nm Fig. 7. TEM micrograph showing dislocation configuration under rolling condition in steels investigated: (a) steel A and (b) steel B. (c) TEM image of steel B shows the interaction between nano oxide precipitation and dislocation, which is indicated by arrows.
Table 2 Longitudinal mechanical properties of steels investigated.
Steel A Steel B
Yield strength (MPa)
Tensile strength (MPa)
Elongation (%)
272 711
372 810
33 22
melt directly. Relevant information on the materials fabrication can be found elsewhere [13]. 4.3. Strengthening mechanism The relationship between yield strength and microstructure can be predicted by the following equation [14]:
ss ¼ s0 þ sg þ sp þ sss þ sd Fig. 8. Tensile stress–strain curves for steel B.
no concentration supply, the further growth of the nano particle will not meet thermodynamics and dynamics conditions and their size will always remain in a smaller size range. Through this process, we can obtain a large number of nanoscale particles in
ð2Þ
where s0 is the internal lattice strength; sg , sp , sss and sd are the strengthening effects caused by grain refinement, precipitation, solid solution, and dislocation, respectively. For present interest, the hardening from precipitates and grain refinement plays a dominant role in steel B. The other strengthening mechanism cannot be overlooked either.
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More recently, it was observed that nanoparticles can effectively serve as a catalyst for heterogeneous nucleation of metal grains during solidification [15,16] and restrictions of grain growth by the effect of Zener pinning [17–19] during solidification and solid-state heat treatment for the existence of a large number of nanoparticles distributed in the grain boundary and intracrystalline. In this work, the nanoparticles formed in melt can also be the heterogeneous nucleus of the sulfide, thus to optimize the inclusion. Obviously, as shown in Fig. 2b, the optimized inclusion pinning at the grain boundary will restrain the growth of grains and generate grain refinement hardening. All the above comprehensive functions of the nano precipitated phase result in significantly smaller grain sizes and markedly improved yield strength according to the Hall–Petch relationship. As shown in Fig. 3, the average grain size of steel B is obviously smaller than that of steel A. As described by the Hall–Petch equation [20]:
s ¼ K y d 1=2
ð3Þ
where d is the average ferrite grain size; Ky is a constant, 15:1–18:1 N mm 3=2 . The average ferrite grain size has been estimated using the Image-Pro Plus software. For the steel with an average grain size of 6 μm, a grain refinement hardening effect of about 225 MPa can be expected. The dominant factor for precipitation hardening is generated by a large number of dispersing nanometer particles for this new steel B when compared with steel A. Precipitation hardening lies in the interaction between the second precipitation and the dislocations, causing the blocking of the dislocation movement, contributing to the high tensile strength due to the Orowan hardening mechanism [21,22]. In Fig. 7a, b, a significant change in the microstructure of the steel occurs. Compared with steel A, a remarkable increase in the dislocation density can be observed, suggesting that the finely dispersed precipitates are beneficial for the increase in dislocation density. Furthermore, it is obvious to see that most of the precipitates are tangled with dislocations, as shown in Fig. 7(c). The HRTEM of precipitated Ti3O5 nano-particles indicates that the lattice contiguity between the oxide particles and matrix is evident. While we substantially improve its strength, the coherent or semi-coherent relationship guarantees that when the moving dislocations meet these coherent particles, part of them will pile up and others will continue to keep slipping along the coherent interface [23]. And this will not lead to the aggregation and hinder of dislocations on the grain boundaries; thus we can improve the strength without sacrificing its ductility too much. According to Gladman's theory, using the modified Ashby– Orowan model to calculate the precipitation hardening effect [24,25]: Rp ¼
10μb r ϕ1=2 ln b 5:72r π 3=2
ð4Þ
where μ is the shearing factor (as for ferrite, μ ¼ 80.26 103 MPa); b is the Burgers Vector, 2.48 10–4 μm; r is the radius of the particle;ϕ is the volume fraction of the precipitates. The precipitate radius (r) and the volume fraction of the precipitates (ϕ) have been estimated using small-angle X-ray-scattering (SAXS) and Fit2d software, which are about 6.7 nm and 0.354%, respectively. Consequently, precipitation hardening generated by nano precipitates is determined using Eq. (4) as 183 MPa. Once again, the above value reveals that the finely dispersed precipitates which are formed in melt can not only have a positive effect on precipitation hardening, but also can act as heterogeneous nucleation core thus to refine as-cast structure grains and optimize inclusions.
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5. Conclusions A novel route of fabricating in-situ nano-particles strengthened ferritic steel was investigated in this study. The main conclusions can be drawn as follows: (1) A ferritic steel reinforced with in-situ nanoparticles which have mean diameter of approximately 5 nm and sizes smaller than 10 nm has been fabricated by a novel route. (2) Size of the inclusions which are nucleated on the preferentially formed nano-particles is refined greatly and their distribution is under good control. The inclusion pinning on the grain boundary will restrain the growth of grains. (3) The nano-particles in the investigated steel are Ti3O5. These particles share a coherent or semi-coherent relationship with the matrix. Thus we can improve its strength substantially without sacrificing its plasticity and toughness too much. (4) The contribution of this novel fabricating route to yield strength reaches 439 MPa in the high strength steel. Precipitation of nano-particles is the dominant factor responsible for yield strength increasing.
Acknowledgments The authors are greatly indebted to Prof. Zidong Wang and Dr. Xiaohua Chen for valuable discussions. We also thank Dr Longfei Zuo for his participation of part of the experiments. This work was supported by the Specialized Research Fund for the Doctoral Program of Higher Education of China (No. 20100006110026). References [1] K. Ehrlich, Fusion Eng. Des. 56 (2001) 71–82. [2] J.H. Schneibel, S. Shim, Mater. Sci. Eng. A 488 (2008) 134–138. [3] J.H. Schneibel, C.T. Liu, D.T. Hoelzer, M.J. Mills, P. Sarosi, T. Hayashi, U. Wendt, H. Heyse, Scr. Mater. 57 (2007) 1040–1043. [4] S. Yamashita, S. Watanabe, S. Ohnuki, H. Takahashi, N. Akasaka, S. Ukai, J. Nucl. Mater. 283 (2000) 647–651. [5] L.S. Zhang, S. Ukai, T. Hoshino, S. Hayashi, X.H. Qu, Acta Mater. 57 (2009) 3671–3682. [6] Z.Z. Liu, Y.S. Sang, Steelmaking 23 (2007) 3. [7] W. Yang, X.G. Yang, L.F. Zhang, F.X. Liu, Steelmaking 29 (2013) 6. [8] Y.M. Li, F.X. Zhu, F.P. Cui, J. Northeast. Univ. 28 (2007) 1002–1005. [9] G. Domizzi, G. Anteri, J. Ovejero-Garcia, Corros. Sci. 43 (2001) 325–339. [10] J.I. Takamura, S. Mizoguchi, Metallurgy of oxides in steels. I. Roles of oxides in steels performance[C]//IISC, in: Proceedings of The Sixth International Iron and Steel Congress vol. 1 (1990), pp. 591–597. [11] T. Sawai, M. Wakoh, S. Mizoguchi, J. Iron Steel Inst. Jpn. 82 (7) (1996) 587–592. [12] H. Goto, K. Miyazawa, K. Amaguchi, et al., ISIJ Int. 34 (1994) 414–419. [13] Z.D. Wang, X.W. Wang, Q.S. Wang, Nanotechnology 20 (2009) 075605. [14] F.B. Pickering, Physical metallurgy and the design of steels[M], Applied Science Publication Limited, London, 1978, p. 63. [15] M.P. De Cicco, L.S. Turng, X.C. Li, J.H. Perepezko, Metall. Mater. Trans. A 42 (2011) 2323–2330. [16] A. Belyakov, F.G. Wei, K. Tsuzaki, Y. Mishima, Mater. Sci. Eng. A 471 (2007) 50. [17] E. Nes, N. Ryum, O. Hunderi, Acta Mater. 33 (1985) 11–22. [18] E. Nes, N. Ryum, O. Hunderi, Acta Metall. 33 (1985) 1. [19] U.F. Kocks, H. Meching, Acta Metall. 29 (1981) 1865. [20] X.P. Mao, X.D. Huo, X.J. Sun, Y.Z. Chai, J. Mater. Process. Technol. 210 (2010) 1660–1666. [21] Y.R. Wen, A. Hirata, Z.W. Zhang, T. Fujita, C.T. Liu, J.H. Jiang, M.W. Chen, Acta Mater. 61 (2013) 2133–2147. [22] W.D. Callister, Materials Science and Engineering: An Introduction, Wiley, New York, 2007. [23] C.R. Dong, H.P. Reng, T.Z. Jing, Microalloyed Non Quenched and Tempered Steel, Metallurgical Industry Press, Beijing, 2000. [24] Y.L. Kang, H. Yu, Steel 38 (2003) 20–26. [25] T. Gladman, Miner. Met. Mater. Soc. 3 (1992) 14.