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Dualism of precipitation morphology in high strength low alloy steel
Author's Accepted Manuscript
Dualism of precipitation morphology in High strength low alloy steel Chih-YuanChen RenYang
,
Chien-ChonChen
,
Jer-
www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(14)01529-9 http://dx.doi.org/10.1016/j.msea.2014.12.029 MSA31846
To appear in:
Materials Science & Engineering A
Received date: 27 September 2014 Revised date: 2 December 2014 Accepted date: 8 December 2014 Cite this article as: Chih-YuanChen , Chien-ChonChen , Jer-RenYang , Dualism of precipitation morphology in High strength low alloy steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2014.12.029 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Dualism of Precipitation Morphology in High Strength Low Alloy Steel a, b
a
b
Chih-Yuan, Chen*, a Chien-Chon, Chen, bJer-Ren, Yang*
Department of Energy Engineering, National United University, Miaoli 36003, Taiwan
Department of Materials Science and Engineering, National Taiwan University, Taipei 10617, Taiwan
Abstract While the role of microalloying elements on precipitation strengthening in ferrite matrix during austenite/ferrite transformation is quite clear, some uncertainty still exists concerning the variability of the microhardness distribution of ferrite grains in the isothermal holding condition. The objective of the present study was to clarify the intrinsic characteristics of carbide precipitation morphology in the ferrite matrix under different processing temperatures and times and to correlate it with austenite decomposition kinetics to elucidate why a large microhardness distribution occurs at low isothermal holding temperature. Better understanding of carbide precipitation behavior can help researchers to determine the root cause of variation in microhardness distribution, which would allow metallurgists to produce high quality steels. Measurement with a Vickers hardness indenter revealed that, in specimens isothermally held at 625ºC, the range of Vickers hardness distribution was 240 to 420 after 5 min of isothermal holding, and 270 to 340 after 60 min. For specimens isothermally held at 725ºC, the range of Vickers hardness distribution was 200 to 330 for 5 min of isothermal holding, and 200 to 250 for 60 min. Therefore, the average microhardness decreased with the isothermal holding temperature and time, and a larger range of distribution occurred with short isothermal holding times. Transmission electron microscopy (TEM) images showed that interface precipitation and random precipitation can occur within the same ferrite grain. The reason is that the austenite decomposition
rate varies with transformation temperature and time. An excessively fast austenite/ferrite interface movement velocity, which usually happens in small ferrite grains, would cause these ferrite grains with microalloying elements to exceed their solubility. Furthermore, these microalloying elements will be precipitated randomly after isothermal holding at longer times. Consequently, a large microhardness distribution can usually be detected in specimens with tiny ferrites because some ferrite grains are in a fresh state, without carbides, due to high austenite/ferrite interface movement velocities. Furthermore, one important technological limit that should be kept in mind is the difficulty of developing only one type of precipitation morphology (i.e., interface precipitation or random precipitation) within every ferrite grain.
Keywords: precipitation, Ti-Mo, nano-sized carbide. Corresponding author: Chih-Yuan, Chen, Phone: 886-37-382383 Fax: 886-37-382339, E-mail: [email protected] Jer-Ren, Yang, Phone: 886-2-23620601 Fax: 886-2-23634562, E-mail: [email protected]
1. Introduction There is currently a huge demand for the automobile industry to develop high strength steel without sacrificing formability. A new generation of high strength automobile wheel steels characterized by ferrite microstructure and nano-sized TiMoC precipitates was successfully developed in 2004 [1]. This advanced steel features both ultra-high strength and excellent stretch flange formability due to its uniform microstructure. Since then, many metallurgists have focused on the precipitation of nano-sized carbides, nitrides, copper and intermetallic compounds, such as NiAl, in the ferrite matrix [2-4]. Early results led most researchers to believe that the main or sole mechanism to obtain these tiny size carbides in HSLA steel was interface precipitation [5-7]. However, in recent works, some authors have noted another precipitation morphology (random homogeneous precipitation) also occurring in HSLA steel. For example, Chen et al. used Ti-V bearing steel to find both interface precipitation and random precipitation within a ferrite grain [8]. Yang et al. used a nano-indenter on the Ti-Nb bearing samples and proposed a better dislocation blocking effect for the interface precipitation than for the randomly dispersed precipitation [9]. On the other hand, Chen et al., using a mathematical model and computer simulation, suggested that random array carbides distributed as finely as in interface precipitation would achieve a better strengthening effect [10]. However, it is not known whether it is possible to obtain only this powerful precipitation morphology within every ferrite grain matrix. The objective of the present study was to systematically study the effect of isothermal holding temperature on the precipitation morphology within the ferrite matrix. To accomplish the object, the Vickers hardness distribution of ferrite grains and precipitate morphology were studied. Furthermore, the effect of austenite decomposition kinetics was also examined.
2. Experimental Procedure
The
chemical
composition
of
the
steel
studied
was
Fe-0.06C-
1.5Mn-0.1Si-0.2Ti-0.2Mo-0.004N-0.002S (wt%). The as-received material was HSLA steel plate (with a thickness of about 45 mm) produced through high temperature soaking, hot rolling, and accelerated cooling. In the present work, all isothermal holding treatments were carried out on a Dilatromic III RDP deformation dilatometer produced by Theta Industries, Inc. Before the preparation of dilatometry specimens, the pieces of steel were homogenized at 1250ºC for 2 days while sealed in a quartz tube containing argon and subsequently quenched in water. After the decarburization layer was removed, the specimens were then machined into cylinders 3 mm in diameter and 6 mm in length. After being austenitized at 1200ºC for 3 min, the specimens were directly cooled to isothermal aging temperatures (625ºC and 725ºC) at a cooling rate of about 100ºC/s to prevent carbide precipitating prior to isothermal aging, after which they were isothermally held at those temperatures for 5 min and 60 min respectively before being quenched to room temperature. The samples were mainly characterized by optical microscopy (OM), transmission electron microscopy (TEM, JEM-2000), and field emission gun transmission electron microscopy (FEG-TEM Tecnai F30) equipped with a nanometer probe energy dispersive spectrometer (EDS). The sizes and volume fractions of ferrite grains were estimated in terms of mean intercept length and standard systematic point count method [11, 12]. A micro-hardness measurement of the specimens to be examined optically was taken using a Vickers hardness tester with a load of 100 g. In order to determine the precipitation status in the ferrite, for each steel and processing condition, measurements from 100 indentation tests were used to plot the final microhardness distribution. In order to avoid the influence of strain fields caused by each indentation test and interference with other phases, the position and the distance of each indentation mark were carefully controlled to ensure that the indentation marks were within a ferrite grain. Thin foil specimens were prepared for transmission electron microscopy from 0.25 mm thick discs sliced from the specimens used in the dilatometry experiments. The discs were thinned to 0.5 mm by
abrasion on silicon carbide paper and then electropolished in a twin jet electropolisher using a solution of 5 vol.% perchloric acid+ 25 vol.% glycerol + 70 vol.% ethanol at -2 ºC and 30 V potential.
3. Results and Discussion Fig. 1 (a)-(d) reveals the typical indentation morphology examined by Vickers hardness indenter. The microstructure of a specimen isothermally held at 625ºC, even for just 5 min, is fully ferrite, while the microstructure of the specimen isothermally held at 725ºC is composed of ferrite and martensite. This typical optical microstructure obeys the diffusion transformation characteristic. Namely, the quantity of newly formed ferrite is a function of the isothermal holding temperature and time. As compared to other isothermal holding temperatures, isothermal holding at 625ºC allows transformation to be completed very quickly due to the large
/
!
transformation driving force combined with the comparably high diffusivity of carbon atoms. The numbers marked around the indentation morphology in the figure represent the sequence of each indentation test. The appearance of each indentation reflects the hardness of each test ferrite grain; that is, a harder ferrite grain usually causes a smaller indentation. Therefore, each indentation mark truly reflects the hardness of each ferrite grain. After 100 indentation tests were completed for every experimental condition, the Vickers hardness distribution was plotted, and the plots are shown in fig. 1(e)-(f). For the specimen isothermally held at 625ºC, the range of Vickers hardness distribution was 240 to 420 for isothermal holding times of 5 min, and 270 to 340 for 60 min. For the specimen isothermally held at 725ºC, the range of Vickers hardness distribution was 200 to 330 for isothermal holding times of 5 min, and 200 to 250 for 60 min. Therefore, it is worthy to note that regardless of the isothermal holding temperature, greater scattering of Vickers hardness occurred with the shorter isothermal holding time, which implies the larger variability in the mechanical strength of each ferrite grain. Nevertheless, the narrowing of the Vickers hardness distribution with a longer isothermal holding time reveals that the mechanical properties of each
ferrite grain can be more uniform when isothermal holding is longer. Furthermore, the average Vickers hardness increased as the isothermal holding temperature decreased, rising from Hv 235 (725ºC, 60 min) to Hv 321 (625ºC, 60 min). It is well known that the hardening contribution to the overall strength of ferrite containing precipitation carbides results from the sum of different hardening mechanisms, including friction stress, solid solution strengthening, grain size refinement strengthening, dislocation strengthening, and precipitation strengthening. In many studies employing similar steel compositions and processing conditions [1, 2, 13, 14], almost every result emphasizes that the main strengthening effect occurring in the HSLA steel is the precipitation strengthening mechanism. Therefore, it is believed that the Vickers hardness distribution of ferrite grains can be mainly ascribed to the precipitation status (ex: size, precipitation morphology, and volume density) within the ferrite grains. This will be explained further in the following section. Table 1 shows variations in average ferrite grain size and volume fractions with different isothermal holding times for the present steel. Fully ferrite grains with smaller sizes could be obtained with isothermal holding at lower temperature. Accompanying the above hardness measurement results, a very interesting trend can be put forward as follows: The smaller the ferrite grain size, the greater the microhardness and the larger the variability in the microhardness distribution. The profound influence of ferrite grain size on the hardness, especially its distribution, can be correlated with austenite decomposition kinetics. It is generally accepted that smaller ferrite grains are usually associated with lower processing temperature, which favors the acceleration of austenite decomposition transformation due to the larger austenite/ferrite transformation driving force. It is thus that some ferrite grains form without interface precipitation as a result of the higher austenite/ferrite interface movement velocity. Therefore the carbide’s non-uniform precipitation status occurring in the ferrite grains would result in larger variability of microhardness distribution in the initial isothermal holding period.
As shown in the fig. 1 (c)-(d), the larger ferrite grains nucleate along the prior austenite grain boundary and grow more slowly into the interior of the parent grain. Therefore, it is believed that the dominant precipitation morphology occurring in the high temperature formation of ferrite grains is primarily interface precipitation due to a lower interface movement velocity. However, based on the analysis of the above austenite decomposition kinetics, the carbide precipitation mechanism can be changed when ferrite grain size is reduced, which usually indicates that the austenite /ferrite interface movement velocity has also altered. Therefore, for smaller ferrite grains formed by hot deformation, the carbide precipitation mechanism is not interface precipitation but random precipitation, since the high austenite/ferrite interface movement velocity results in supersaturated ferrite grains. Fig. 2 shows bright field and dark field TEM images and corresponding diffraction patterns of steel isothermally held at 625ºC for 60 min. Interestingly, three different precipitation morphologies (zone A, zone B, and zone C) were observed in the same ferrite grain. In zone A, the precipitation carbides exhibited a complex precipitation morphology with sizes of about 6 to 8 nm. Very few carbides precipitated as a regular array (i.e., interface precipitation), and most carbides existed as a random distribution status (i.e., random precipitation). Carbides with a regular array appearance (i.e., interface precipitation) were found in zone B, despite the larger particle size (13 to 16 nm). On the other hand, the precipitated MC carbides with an average size of 4 to 6 nm showed a randomly and homogeneously distributed morphology (i.e., random precipitation) in zone C. The corresponding diffraction patterns of zone A and B exhibited only one variant of the Baker-Nutting (BN) relationship between the precipitation carbide and the ferrite matrix. It is generally accepted that the MC carbide has a NaCl-type crystal structure and obeys the BN orientation relationship with respect to the ferrite matrix [15]: (001)carbide
(001)ferrite
[110]carbide [100]ferrite.
In this orientation relationship, the {100} planes of carbides and {100} planes of ferrite are coherent, so carbides grow along the habit planes of {100} ferrite, leading to a disc-shaped morphology of the carbides. However, in the case of interface precipitation, unlike with other variants, the variant with the smallest angle between the interface of ferrite and the MC carbide broad plane prefers to grow smoothly along this direction. The coexistence of different precipitation morphologies in the ferrite matrix can be ascribed to variation in the austenite decomposition rate. Interface precipitation of carbides usually occurs within large ferrite grains due to a suitable gamma/alpha interface movement velocity. On the other hand, smaller ferrite grains can usually be obtained with isothermal holding at low temperature, meaning fewer microalloying elements have the opportunity to be precipitated in the interface distribution mode due to the fast decomposition of austenite at low temperature. In other words, supersaturated ferrite grains usually occur when isothermal holding temperature is low. Despite the existence of a few interface precipitation carbides within the supersaturated ferrite matrix, microalloying elements exceeding the solubility limit can still be precipitated with a random distributed morphology within the supersaturated ferrite grains. Though very tiny in size, the carbides in fig. 2 still exhibit great variability in size. This phenomenon can be mainly ascribed to the points at which different carbide precipitation mechanisms begin to affect the transformation. Carbides with interface precipitation morphology occurring in the early period of transformation have more time to grow and become relatively large in size. Random precipitated carbides, on the other hand, usually appear in the later stage and thus have little time to grow further. Moreover, the Ashby-Orowan equation [16] suggests that these nanometer carbides can contribute greatly to the strength. The coexistence of interface precipitation and random precipitation in the same ferrite grain can be mainly ascribed to the characteristics of diffusion transformation and will be explained in greater detail in the following section.
In order to evaluate the effect of precipitation strengthening, specimens that were isothermally held only at low temperature were tested because a fully ferrite microstructure, instead of two phases (i.e., ferrite and martensite) occurs only at low temperature. Table 2 presents the mechanical properties of a specimen isothermally held at low temperature. For ferritic steels, the contributed yield strength due to the friction stress of the ferritic matrix, solid solution strengthening and grain boundary strengthening can be expressed as [17]:
1.13 ] (1) D is the increased yield stress in MPa
Cıy =Cσ 0 + Cσ ss +Cσ gb = 15.4[3.5 + 2.1( Mn ) + 5.4 ( Si ) + 23.4 ( C ) + 23 ( Nf ) + where Cσ 0 is the friction stress of pure iron in MPa, Cσ ss
due to solid solute strengthening, Cσ gb is the increased yield stress in MPa due to grain boundary strengthening, (Mn), (Si), (C), and (N) are the average wt% in solution. D is ferrite grain diameter in mm. The (Mn) and (Si) in ferrite matrix are 1.5 wt% and 0.1 wt%, respectively. The carbon content in ferrite matrix should be below 0.02 wt%, and the free nitrogen content approaches 0 because Ti was added to the steel to fix nitrogen by the formation of stable TiN. Furthermore, the dislocation strengthening can be estimated as [16]:
Cσ dis = M α Gb ρ (2) where M is a Taylor factor of 2.75, Į is a constant of ~0.435 [16], G is a shear modulus of 80.3 GPa, b is a Burgers vector of 0.248nm, and ȡ is dislocation density for ~ 5 × 10 −13 m −2 in the present study. The total yield strength can be calculated with the following equation [8]: 2 2 ıy =Cσ 0 +Cσ ss +Cσ gb + (Cσ )dis + (Cσ )Orowan
(3)
The components of the calculated yield stress by Eqs. (1) - (3) and the total yield stress are shown in table 3. The major decreases in yield stress with increasing isothermal holding time result from the growth of carbides, lowering the contribution from Orowan strengthening. Fig. 3 illustrates the traditional TTT curve for diffusional phase transformation, the corresponding transformation fraction figure, and the evolution of precipitation carbides within ferrite grains accompanying the austenite/ferrite transformation process. The TTT curve and corresponding transformation fraction figure can also be found elsewhere [18]. Based on this figure, the transformation of austenite to ferrite can be divided into three individual stages
according to transformation rate. In the initial transformation stage, polygonal ferrite grains usually nucleate at prior austenite grain boundaries and grow steadily along the grain boundary. In this period, interface precipitation carbides can be found in most ferrite grains because the interface movement velocity is suitable for heterogeneous nucleation of microalloying elements at the austenite/ferrite interface (i.e., microalloying elements, ledge, semi-ledge, step plane, and terrace plane). As shown in fig. 3, evidence of a lower interface movement velocity can be found in the traditional TTT curve. However, in the intermediate stage of transformation, the interface movement velocity is higher than in the other transformation periods, so it is reasonable to assume that the diffusivity of microalloying elements is not comparable to the interface movement velocity, and thus most microalloying elements still remain in the ferrite grains. Thus, these remaining microalloying elements can be randomly and homogeneously precipitated spontaneously inside the ferrite grains due to the large difference in solubility between austenite and ferrite [16]. Although random homogeneous precipitation can be commonly obtained in this stage, the interface precipitation morphology still has a chance of occurring in spite of the high interface movement velocity. The reason could be that the movement at the austenite/ferrite interface is hindered by solute atoms, and thus the movement velocity is reduced. Consequently, it is common to observe two precipitation morphologies coexisting within one ferrite grain. In the final stage of transformation, the austenite decomposition again slows, and thus carbides are expected to nucleate along the austenite/ferrite interface in an ordered array. The large scattering of Vickers hardness distributions with a short isothermal holding time can be associated with the quantity of carbide precipitates in the ferrite matrix. It reasonable to suggest that harder grains that formed in the initial stage of transformation would contain interface precipitation carbides within the ferrite matrix and that softer ferrite grains that formed in the early period of the intermediate stage of transformation would exist in a fresh state, containing no precipitation particles, due to the fast interface movement velocity. Afterwards, the randomly and homogeneously precipitated carbides would nucleate and grow within these fresh
ferrite grains during the longer isothermal times. In this stage, the range of Vickers hardness distribution decreases, as most ferrite grains can be strengthened effectively by nano-sized carbides. In summary, whether interface precipitation or random homogeneous precipitation occurs is primarily determined by the austenite/ferrite interface movement velocity and the diffusivity of microalloying elements, which are mainly influenced by the isothermal holding temperature. Therefore, random homogeneous precipitation is obtained in most ferrite grains, as the transformation can be completed quickly at a low isothermal holding temperature.
4. Conclusion 1. At low- and high-isothermal holding temperatures, two different carbide precipitation morphologies, namely interface precipitation and random homogeneous precipitation, can occur within the same ferrite grain. This dualism of the precipitation morphology phenomenon is a characteristic of diffusional phase transformation, since the influential factor determining precipitation morphology mainly depends on the austenite/ferrite movement velocity, which is affected by the isothermal holding temperature and time. The diffusional phase transformation rate is not constant, usually increasing in speed in the intermediate stage of transformation. Therefore, it is difficult to obtain only one kind of carbide precipitation morphology, interface precipitation, within each ferrite grain. 2. The Vickers hardness distribution of ferrite grains in Ti-Mo bearing steel was more scattered after short isothermal holding times than after longer times. The reason is that some softer ferrite grains were in a “fresh” state and were strengthened by randomly and homogeneously dispersed carbides during the longer isothermal holding times. 3. The lower average Vickers hardness in samples held at the higher isothermal holding temperature is related to the larger size of the precipitated carbides. Therefore, the most
important factor in the final strength of ferrite steel is the isothermal holding temperature, not the isothermal holding time. 4. Although it is very difficult to obtain only random precipitation carbides within every ferrite grain (i.e., a technological limit), it is possible to obtain only random precipitated carbides within most ferrite grains if the austenite/ferrite transformation is accelerated. For example, despite isothermal holding at high temperature, heavily deformed steels can have small ferrite grains with less variability in microhardness distribution due to acceleration of the decomposition of austenite.
References 1. Yoshimasa Funakawa, Tsuyoshi Shiozaki, Kunikazu Tomita, Tetsuo Yamamoto, and Eiji Maeda, Development of High Strength Hot-rolled Sheet Steel Consisting of Ferrite and Nanometer-sized Carbides, ISIJ inter., 2004; 11: 1945-51. 2. C.Y. Chen, H.W. Yen, F.H. Kao, W.C. Li, C.Y. Huang, J. R. Yang and S.H.Wang, Precipitation hardening of high-strength low-alloy steels by nanometer-sized carbides, Mater. Sci. Eng. A, 2009; 499: 162-66. 3. Z.B. Jiao, J.H. Luan, Z.W. Zhang, M.K. Miller, W.B. Ma, and C.T. Liu, Synergistic effects of Cu and Ni on nanoscale precipitation and mechanical properties of high-strength steels, Acta Mater., 2013; 61: 5996-6055. 4. Y.R. Wen, Y.P. Li, A. Hirata, Y. Zhang, T. Fujita, T. Furuhara, C.T. Liu, A. Chiba, M.W. Chen, Synergistic alloying effect on microstructural evolution and mechanical properties of Cu precipitation-strengthened ferritic alloys, Acta Mater., 2013; 61: 7726-7740. 5. A. T. Davenport, F. G. Berry, and R. W. K. Honeycombe, Interface Precipitation in iron alloys, Met. Sci., 1968, 2: 104-104. 6. R. W. K. Honeycombe, Transformation from Austenite in Alloy Steels, Metall. Trans. A, 1976; 7A: 915-936. 7. Jae Hoon Jang, Chang-Hoon Lee, Yoon-Uk Heo, and Dong-Woo Suh, Stability of (Ti, M)C (M = Nb, V, Mo and W) carbide in steels using first-principles calculations, Acta Mater. 2012; 60: 208-217. 8. Jun Chen, Meng-yang Lv, Shuai Tang, Zhen-yu Liu, Guo-dong Wang, Influence of cooling paths on microstructural characteristics and precipitation behaviors in a low carbon V–Ti
microalloyed steel. Mater. Sci. Eng. A, 2014; 594: 389-393. 9. Yang Xu, Weina Zhang, Mingxue Sun, Hailong Yi, Zhenyu Liu, The blocking effects of interface precipitation on dislocations’ movement in Ti-bearing micro-alloyed steel, Mater. Let., 2015; 139: 177-181. 10. M.-Y. Chen, M. Gouné, M. Verdier, Y. Bréchet, J.-R. Yang, Interphase precipitation in vanadium-alloyed steels: Strengthening contribution and morphological variability with austenite to ferrite transformation, Acta Mater., 2014; 64: 78-92. 11. A. Mein, G. Fourlaris, D. Crowther, and P.J. Evans: ‘The influence of aluminium on the ferrite formation and microstructural development in hot rolled dual-phase steel l’, Mater. Charact., 64 (2012) 69-78. 12. S. Shanmugam, R.D.K. Misra, T. Mannering, D. Panda, S.G. Jansto: ‘Impact toughness and microstructure relationship in niobium- and vanadium-microalloyed steels processed with varied cooling rates to similar yield strength’, Mater. Sci. Eng. A 437 (2006) 436-445. 13. Dae-Bum Park, Moo-Young Huh, Jae-Hyeok Shim, Jin-Yoo Suh, Kyu-Ho Lee, Woo-Sang Jung, Strengthening mechanism of hot rolled Ti and Nb microalloyed HSLA steels containing Mo and W with various coiling temperature, Mater. Sci. Eng. A, 2013; 560: 528-534. 14. Yong Woo Kim, Seok Weon Song, Seok Jong Seo, Seong-Gu Hong, Chong Soo Lee, Development of Ti and Mo micro-alloyed hot-rolled high strength sheet steel by controlling thermomechanical controlled processing schedule, Mater. Sci. Eng. A, 2013; 565: 430-438. 15. Y.-J. Zhang, G. Miyamoto, K. Shinbo, T. Furuhara, Effects of Į/Ȗ orientation relationship on VC interphase precipitation in low-carbon steels, Scripta Mater., 2013; 69: 17-20. 16. T, Gladman, Precipitation hardening in metals, Mater. Sci, Technol., 1999, 15: 30-36. 17. F. B. Pickering, Physical Metallurgy and the Design of Steels, Applied Science Publishing Ltd., London, 1978, 63. 18. .David A. Porter, and Kenneth E. Easterling, Phase Transformations in Metals and Alloys, 2rd ed. Chapman & Hall press, 1992.
Fig. 1 Optical micrographs and corresponding Vickers hardness distribution of Ti-Mo steel after different heat treatments: (a, e) isothermal holding at 625ºC for 5 min, (b, f) isothermal holding at 625ºC for 60 min, (c, g) isothermal holding at 725ºC for 5 min, (d, h) isothermal holding at 725ºC for 60 min. The number around each indentation represents the indentation test sequence. It is worthy to note that the size of the indentation does not change during the test sequence, which implies that the size of later indentations was not influenced by strain fields originating from previous indents. Fig. 2 TEM images of Ti-Mo bearing steel after isothermal holding at 625ºC for 60 min; (a) low- and (b, e, h) high-magnification TEM bright-field images showing two different precipitation morphologies (i.e., interface precipitation (zone A, zone B) and random homogeneous precipitation (zone C)) within the same polygonal ferrite grain, (c, f, i) high magnification dark field TEM images (using (200)TiMoC reflection of precipitation carbides, revealing the morphology of precipitates), and the (d, g, j) corresponding diffraction patterns, revealing the BN orientation relationship between precipitated carbides and ferrite matrix. Fig. 3 Schematic drawing illustrating the influence of isothermal holding temperature on the precipitation morphology within the ferrite matrix. Coexistence of different precipitation morphologies, namely interface precipitation and random homogeneous precipitation, can be observed in the HSLA steel that has undergone the general manufacturing process.
Figure1a-d
Figure1e-h
Figure2
Figure3
Table1
Table1 The microstructure characterization of Ti-Mo bearing steel after isothermal treatment at 625ºC and 725ºC for 5 and 60 min. Ti-Mo bearing steel presents all ferrite phase in the low isothermal holding temperature after very shortly isothermal holding times (5 min), indicating acceleration of the decomposition of austenite.
625кʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ˊ˅ˈк Isothermal holding time Ferrite Area fraction (%) Grain size (μm)
5 min 100 23.6̈́1.6
60 min 100 32.3̈́1.8
5 min
60 min
12.8̈́˄ˁ˅ʳ ʳ ʳ 28.7̈́˄ˁˉʳ ʳ .21.5̈́1.2
41.2̈́2.7
Table2
Table 2 Mechanical properties of Ti-Mo bearing steel after isothermal treatment at 625ºC for 5 and 60 min.
Process
Ys (MPa)
Ts (MPa)
A(%)
YR
625ʚ_5min
625ʚ_60min
Table3
Table 3 Components of yield stress in the steels studied.
Process
625
_5min
625 _60min
οߪ(MPa)
53.9
53.9
οߪ௦௦ (MPa)
.64.1
64.1
ο(MPa)
113.0
96.8
οߪௗ௦ (MPa)
168.0
168.0
οߪ௪ (MPa)
.274.0
243.2
ߪ௬ (MPa)
.673.0
626.0
Dualism of precipitation morphology in High strength low alloy steel Chih-YuanChen RenYang
,
Chien-ChonChen
,
Jer-
www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(14)01529-9 http://dx.doi.org/10.1016/j.msea.2014.12.029 MSA31846
To appear in:
Materials Science & Engineering A
Received date: 27 September 2014 Revised date: 2 December 2014 Accepted date: 8 December 2014 Cite this article as: Chih-YuanChen , Chien-ChonChen , Jer-RenYang , Dualism of precipitation morphology in High strength low alloy steel, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2014.12.029 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Dualism of Precipitation Morphology in High Strength Low Alloy Steel a, b
a
b
Chih-Yuan, Chen*, a Chien-Chon, Chen, bJer-Ren, Yang*
Department of Energy Engineering, National United University, Miaoli 36003, Taiwan
Department of Materials Science and Engineering, National Taiwan University, Taipei 10617, Taiwan
Abstract While the role of microalloying elements on precipitation strengthening in ferrite matrix during austenite/ferrite transformation is quite clear, some uncertainty still exists concerning the variability of the microhardness distribution of ferrite grains in the isothermal holding condition. The objective of the present study was to clarify the intrinsic characteristics of carbide precipitation morphology in the ferrite matrix under different processing temperatures and times and to correlate it with austenite decomposition kinetics to elucidate why a large microhardness distribution occurs at low isothermal holding temperature. Better understanding of carbide precipitation behavior can help researchers to determine the root cause of variation in microhardness distribution, which would allow metallurgists to produce high quality steels. Measurement with a Vickers hardness indenter revealed that, in specimens isothermally held at 625ºC, the range of Vickers hardness distribution was 240 to 420 after 5 min of isothermal holding, and 270 to 340 after 60 min. For specimens isothermally held at 725ºC, the range of Vickers hardness distribution was 200 to 330 for 5 min of isothermal holding, and 200 to 250 for 60 min. Therefore, the average microhardness decreased with the isothermal holding temperature and time, and a larger range of distribution occurred with short isothermal holding times. Transmission electron microscopy (TEM) images showed that interface precipitation and random precipitation can occur within the same ferrite grain. The reason is that the austenite decomposition
rate varies with transformation temperature and time. An excessively fast austenite/ferrite interface movement velocity, which usually happens in small ferrite grains, would cause these ferrite grains with microalloying elements to exceed their solubility. Furthermore, these microalloying elements will be precipitated randomly after isothermal holding at longer times. Consequently, a large microhardness distribution can usually be detected in specimens with tiny ferrites because some ferrite grains are in a fresh state, without carbides, due to high austenite/ferrite interface movement velocities. Furthermore, one important technological limit that should be kept in mind is the difficulty of developing only one type of precipitation morphology (i.e., interface precipitation or random precipitation) within every ferrite grain.
Keywords: precipitation, Ti-Mo, nano-sized carbide. Corresponding author: Chih-Yuan, Chen, Phone: 886-37-382383 Fax: 886-37-382339, E-mail: [email protected] Jer-Ren, Yang, Phone: 886-2-23620601 Fax: 886-2-23634562, E-mail: [email protected]
1. Introduction There is currently a huge demand for the automobile industry to develop high strength steel without sacrificing formability. A new generation of high strength automobile wheel steels characterized by ferrite microstructure and nano-sized TiMoC precipitates was successfully developed in 2004 [1]. This advanced steel features both ultra-high strength and excellent stretch flange formability due to its uniform microstructure. Since then, many metallurgists have focused on the precipitation of nano-sized carbides, nitrides, copper and intermetallic compounds, such as NiAl, in the ferrite matrix [2-4]. Early results led most researchers to believe that the main or sole mechanism to obtain these tiny size carbides in HSLA steel was interface precipitation [5-7]. However, in recent works, some authors have noted another precipitation morphology (random homogeneous precipitation) also occurring in HSLA steel. For example, Chen et al. used Ti-V bearing steel to find both interface precipitation and random precipitation within a ferrite grain [8]. Yang et al. used a nano-indenter on the Ti-Nb bearing samples and proposed a better dislocation blocking effect for the interface precipitation than for the randomly dispersed precipitation [9]. On the other hand, Chen et al., using a mathematical model and computer simulation, suggested that random array carbides distributed as finely as in interface precipitation would achieve a better strengthening effect [10]. However, it is not known whether it is possible to obtain only this powerful precipitation morphology within every ferrite grain matrix. The objective of the present study was to systematically study the effect of isothermal holding temperature on the precipitation morphology within the ferrite matrix. To accomplish the object, the Vickers hardness distribution of ferrite grains and precipitate morphology were studied. Furthermore, the effect of austenite decomposition kinetics was also examined.
2. Experimental Procedure
The
chemical
composition
of
the
steel
studied
was
Fe-0.06C-
1.5Mn-0.1Si-0.2Ti-0.2Mo-0.004N-0.002S (wt%). The as-received material was HSLA steel plate (with a thickness of about 45 mm) produced through high temperature soaking, hot rolling, and accelerated cooling. In the present work, all isothermal holding treatments were carried out on a Dilatromic III RDP deformation dilatometer produced by Theta Industries, Inc. Before the preparation of dilatometry specimens, the pieces of steel were homogenized at 1250ºC for 2 days while sealed in a quartz tube containing argon and subsequently quenched in water. After the decarburization layer was removed, the specimens were then machined into cylinders 3 mm in diameter and 6 mm in length. After being austenitized at 1200ºC for 3 min, the specimens were directly cooled to isothermal aging temperatures (625ºC and 725ºC) at a cooling rate of about 100ºC/s to prevent carbide precipitating prior to isothermal aging, after which they were isothermally held at those temperatures for 5 min and 60 min respectively before being quenched to room temperature. The samples were mainly characterized by optical microscopy (OM), transmission electron microscopy (TEM, JEM-2000), and field emission gun transmission electron microscopy (FEG-TEM Tecnai F30) equipped with a nanometer probe energy dispersive spectrometer (EDS). The sizes and volume fractions of ferrite grains were estimated in terms of mean intercept length and standard systematic point count method [11, 12]. A micro-hardness measurement of the specimens to be examined optically was taken using a Vickers hardness tester with a load of 100 g. In order to determine the precipitation status in the ferrite, for each steel and processing condition, measurements from 100 indentation tests were used to plot the final microhardness distribution. In order to avoid the influence of strain fields caused by each indentation test and interference with other phases, the position and the distance of each indentation mark were carefully controlled to ensure that the indentation marks were within a ferrite grain. Thin foil specimens were prepared for transmission electron microscopy from 0.25 mm thick discs sliced from the specimens used in the dilatometry experiments. The discs were thinned to 0.5 mm by
abrasion on silicon carbide paper and then electropolished in a twin jet electropolisher using a solution of 5 vol.% perchloric acid+ 25 vol.% glycerol + 70 vol.% ethanol at -2 ºC and 30 V potential.
3. Results and Discussion Fig. 1 (a)-(d) reveals the typical indentation morphology examined by Vickers hardness indenter. The microstructure of a specimen isothermally held at 625ºC, even for just 5 min, is fully ferrite, while the microstructure of the specimen isothermally held at 725ºC is composed of ferrite and martensite. This typical optical microstructure obeys the diffusion transformation characteristic. Namely, the quantity of newly formed ferrite is a function of the isothermal holding temperature and time. As compared to other isothermal holding temperatures, isothermal holding at 625ºC allows transformation to be completed very quickly due to the large
/
!
transformation driving force combined with the comparably high diffusivity of carbon atoms. The numbers marked around the indentation morphology in the figure represent the sequence of each indentation test. The appearance of each indentation reflects the hardness of each test ferrite grain; that is, a harder ferrite grain usually causes a smaller indentation. Therefore, each indentation mark truly reflects the hardness of each ferrite grain. After 100 indentation tests were completed for every experimental condition, the Vickers hardness distribution was plotted, and the plots are shown in fig. 1(e)-(f). For the specimen isothermally held at 625ºC, the range of Vickers hardness distribution was 240 to 420 for isothermal holding times of 5 min, and 270 to 340 for 60 min. For the specimen isothermally held at 725ºC, the range of Vickers hardness distribution was 200 to 330 for isothermal holding times of 5 min, and 200 to 250 for 60 min. Therefore, it is worthy to note that regardless of the isothermal holding temperature, greater scattering of Vickers hardness occurred with the shorter isothermal holding time, which implies the larger variability in the mechanical strength of each ferrite grain. Nevertheless, the narrowing of the Vickers hardness distribution with a longer isothermal holding time reveals that the mechanical properties of each
ferrite grain can be more uniform when isothermal holding is longer. Furthermore, the average Vickers hardness increased as the isothermal holding temperature decreased, rising from Hv 235 (725ºC, 60 min) to Hv 321 (625ºC, 60 min). It is well known that the hardening contribution to the overall strength of ferrite containing precipitation carbides results from the sum of different hardening mechanisms, including friction stress, solid solution strengthening, grain size refinement strengthening, dislocation strengthening, and precipitation strengthening. In many studies employing similar steel compositions and processing conditions [1, 2, 13, 14], almost every result emphasizes that the main strengthening effect occurring in the HSLA steel is the precipitation strengthening mechanism. Therefore, it is believed that the Vickers hardness distribution of ferrite grains can be mainly ascribed to the precipitation status (ex: size, precipitation morphology, and volume density) within the ferrite grains. This will be explained further in the following section. Table 1 shows variations in average ferrite grain size and volume fractions with different isothermal holding times for the present steel. Fully ferrite grains with smaller sizes could be obtained with isothermal holding at lower temperature. Accompanying the above hardness measurement results, a very interesting trend can be put forward as follows: The smaller the ferrite grain size, the greater the microhardness and the larger the variability in the microhardness distribution. The profound influence of ferrite grain size on the hardness, especially its distribution, can be correlated with austenite decomposition kinetics. It is generally accepted that smaller ferrite grains are usually associated with lower processing temperature, which favors the acceleration of austenite decomposition transformation due to the larger austenite/ferrite transformation driving force. It is thus that some ferrite grains form without interface precipitation as a result of the higher austenite/ferrite interface movement velocity. Therefore the carbide’s non-uniform precipitation status occurring in the ferrite grains would result in larger variability of microhardness distribution in the initial isothermal holding period.
As shown in the fig. 1 (c)-(d), the larger ferrite grains nucleate along the prior austenite grain boundary and grow more slowly into the interior of the parent grain. Therefore, it is believed that the dominant precipitation morphology occurring in the high temperature formation of ferrite grains is primarily interface precipitation due to a lower interface movement velocity. However, based on the analysis of the above austenite decomposition kinetics, the carbide precipitation mechanism can be changed when ferrite grain size is reduced, which usually indicates that the austenite /ferrite interface movement velocity has also altered. Therefore, for smaller ferrite grains formed by hot deformation, the carbide precipitation mechanism is not interface precipitation but random precipitation, since the high austenite/ferrite interface movement velocity results in supersaturated ferrite grains. Fig. 2 shows bright field and dark field TEM images and corresponding diffraction patterns of steel isothermally held at 625ºC for 60 min. Interestingly, three different precipitation morphologies (zone A, zone B, and zone C) were observed in the same ferrite grain. In zone A, the precipitation carbides exhibited a complex precipitation morphology with sizes of about 6 to 8 nm. Very few carbides precipitated as a regular array (i.e., interface precipitation), and most carbides existed as a random distribution status (i.e., random precipitation). Carbides with a regular array appearance (i.e., interface precipitation) were found in zone B, despite the larger particle size (13 to 16 nm). On the other hand, the precipitated MC carbides with an average size of 4 to 6 nm showed a randomly and homogeneously distributed morphology (i.e., random precipitation) in zone C. The corresponding diffraction patterns of zone A and B exhibited only one variant of the Baker-Nutting (BN) relationship between the precipitation carbide and the ferrite matrix. It is generally accepted that the MC carbide has a NaCl-type crystal structure and obeys the BN orientation relationship with respect to the ferrite matrix [15]: (001)carbide
(001)ferrite
[110]carbide [100]ferrite.
In this orientation relationship, the {100} planes of carbides and {100} planes of ferrite are coherent, so carbides grow along the habit planes of {100} ferrite, leading to a disc-shaped morphology of the carbides. However, in the case of interface precipitation, unlike with other variants, the variant with the smallest angle between the interface of ferrite and the MC carbide broad plane prefers to grow smoothly along this direction. The coexistence of different precipitation morphologies in the ferrite matrix can be ascribed to variation in the austenite decomposition rate. Interface precipitation of carbides usually occurs within large ferrite grains due to a suitable gamma/alpha interface movement velocity. On the other hand, smaller ferrite grains can usually be obtained with isothermal holding at low temperature, meaning fewer microalloying elements have the opportunity to be precipitated in the interface distribution mode due to the fast decomposition of austenite at low temperature. In other words, supersaturated ferrite grains usually occur when isothermal holding temperature is low. Despite the existence of a few interface precipitation carbides within the supersaturated ferrite matrix, microalloying elements exceeding the solubility limit can still be precipitated with a random distributed morphology within the supersaturated ferrite grains. Though very tiny in size, the carbides in fig. 2 still exhibit great variability in size. This phenomenon can be mainly ascribed to the points at which different carbide precipitation mechanisms begin to affect the transformation. Carbides with interface precipitation morphology occurring in the early period of transformation have more time to grow and become relatively large in size. Random precipitated carbides, on the other hand, usually appear in the later stage and thus have little time to grow further. Moreover, the Ashby-Orowan equation [16] suggests that these nanometer carbides can contribute greatly to the strength. The coexistence of interface precipitation and random precipitation in the same ferrite grain can be mainly ascribed to the characteristics of diffusion transformation and will be explained in greater detail in the following section.
In order to evaluate the effect of precipitation strengthening, specimens that were isothermally held only at low temperature were tested because a fully ferrite microstructure, instead of two phases (i.e., ferrite and martensite) occurs only at low temperature. Table 2 presents the mechanical properties of a specimen isothermally held at low temperature. For ferritic steels, the contributed yield strength due to the friction stress of the ferritic matrix, solid solution strengthening and grain boundary strengthening can be expressed as [17]:
1.13 ] (1) D is the increased yield stress in MPa
Cıy =Cσ 0 + Cσ ss +Cσ gb = 15.4[3.5 + 2.1( Mn ) + 5.4 ( Si ) + 23.4 ( C ) + 23 ( Nf ) + where Cσ 0 is the friction stress of pure iron in MPa, Cσ ss
due to solid solute strengthening, Cσ gb is the increased yield stress in MPa due to grain boundary strengthening, (Mn), (Si), (C), and (N) are the average wt% in solution. D is ferrite grain diameter in mm. The (Mn) and (Si) in ferrite matrix are 1.5 wt% and 0.1 wt%, respectively. The carbon content in ferrite matrix should be below 0.02 wt%, and the free nitrogen content approaches 0 because Ti was added to the steel to fix nitrogen by the formation of stable TiN. Furthermore, the dislocation strengthening can be estimated as [16]:
Cσ dis = M α Gb ρ (2) where M is a Taylor factor of 2.75, Į is a constant of ~0.435 [16], G is a shear modulus of 80.3 GPa, b is a Burgers vector of 0.248nm, and ȡ is dislocation density for ~ 5 × 10 −13 m −2 in the present study. The total yield strength can be calculated with the following equation [8]: 2 2 ıy =Cσ 0 +Cσ ss +Cσ gb + (Cσ )dis + (Cσ )Orowan
(3)
The components of the calculated yield stress by Eqs. (1) - (3) and the total yield stress are shown in table 3. The major decreases in yield stress with increasing isothermal holding time result from the growth of carbides, lowering the contribution from Orowan strengthening. Fig. 3 illustrates the traditional TTT curve for diffusional phase transformation, the corresponding transformation fraction figure, and the evolution of precipitation carbides within ferrite grains accompanying the austenite/ferrite transformation process. The TTT curve and corresponding transformation fraction figure can also be found elsewhere [18]. Based on this figure, the transformation of austenite to ferrite can be divided into three individual stages
according to transformation rate. In the initial transformation stage, polygonal ferrite grains usually nucleate at prior austenite grain boundaries and grow steadily along the grain boundary. In this period, interface precipitation carbides can be found in most ferrite grains because the interface movement velocity is suitable for heterogeneous nucleation of microalloying elements at the austenite/ferrite interface (i.e., microalloying elements, ledge, semi-ledge, step plane, and terrace plane). As shown in fig. 3, evidence of a lower interface movement velocity can be found in the traditional TTT curve. However, in the intermediate stage of transformation, the interface movement velocity is higher than in the other transformation periods, so it is reasonable to assume that the diffusivity of microalloying elements is not comparable to the interface movement velocity, and thus most microalloying elements still remain in the ferrite grains. Thus, these remaining microalloying elements can be randomly and homogeneously precipitated spontaneously inside the ferrite grains due to the large difference in solubility between austenite and ferrite [16]. Although random homogeneous precipitation can be commonly obtained in this stage, the interface precipitation morphology still has a chance of occurring in spite of the high interface movement velocity. The reason could be that the movement at the austenite/ferrite interface is hindered by solute atoms, and thus the movement velocity is reduced. Consequently, it is common to observe two precipitation morphologies coexisting within one ferrite grain. In the final stage of transformation, the austenite decomposition again slows, and thus carbides are expected to nucleate along the austenite/ferrite interface in an ordered array. The large scattering of Vickers hardness distributions with a short isothermal holding time can be associated with the quantity of carbide precipitates in the ferrite matrix. It reasonable to suggest that harder grains that formed in the initial stage of transformation would contain interface precipitation carbides within the ferrite matrix and that softer ferrite grains that formed in the early period of the intermediate stage of transformation would exist in a fresh state, containing no precipitation particles, due to the fast interface movement velocity. Afterwards, the randomly and homogeneously precipitated carbides would nucleate and grow within these fresh
ferrite grains during the longer isothermal times. In this stage, the range of Vickers hardness distribution decreases, as most ferrite grains can be strengthened effectively by nano-sized carbides. In summary, whether interface precipitation or random homogeneous precipitation occurs is primarily determined by the austenite/ferrite interface movement velocity and the diffusivity of microalloying elements, which are mainly influenced by the isothermal holding temperature. Therefore, random homogeneous precipitation is obtained in most ferrite grains, as the transformation can be completed quickly at a low isothermal holding temperature.
4. Conclusion 1. At low- and high-isothermal holding temperatures, two different carbide precipitation morphologies, namely interface precipitation and random homogeneous precipitation, can occur within the same ferrite grain. This dualism of the precipitation morphology phenomenon is a characteristic of diffusional phase transformation, since the influential factor determining precipitation morphology mainly depends on the austenite/ferrite movement velocity, which is affected by the isothermal holding temperature and time. The diffusional phase transformation rate is not constant, usually increasing in speed in the intermediate stage of transformation. Therefore, it is difficult to obtain only one kind of carbide precipitation morphology, interface precipitation, within each ferrite grain. 2. The Vickers hardness distribution of ferrite grains in Ti-Mo bearing steel was more scattered after short isothermal holding times than after longer times. The reason is that some softer ferrite grains were in a “fresh” state and were strengthened by randomly and homogeneously dispersed carbides during the longer isothermal holding times. 3. The lower average Vickers hardness in samples held at the higher isothermal holding temperature is related to the larger size of the precipitated carbides. Therefore, the most
important factor in the final strength of ferrite steel is the isothermal holding temperature, not the isothermal holding time. 4. Although it is very difficult to obtain only random precipitation carbides within every ferrite grain (i.e., a technological limit), it is possible to obtain only random precipitated carbides within most ferrite grains if the austenite/ferrite transformation is accelerated. For example, despite isothermal holding at high temperature, heavily deformed steels can have small ferrite grains with less variability in microhardness distribution due to acceleration of the decomposition of austenite.
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microalloyed steel. Mater. Sci. Eng. A, 2014; 594: 389-393. 9. Yang Xu, Weina Zhang, Mingxue Sun, Hailong Yi, Zhenyu Liu, The blocking effects of interface precipitation on dislocations’ movement in Ti-bearing micro-alloyed steel, Mater. Let., 2015; 139: 177-181. 10. M.-Y. Chen, M. Gouné, M. Verdier, Y. Bréchet, J.-R. Yang, Interphase precipitation in vanadium-alloyed steels: Strengthening contribution and morphological variability with austenite to ferrite transformation, Acta Mater., 2014; 64: 78-92. 11. A. Mein, G. Fourlaris, D. Crowther, and P.J. Evans: ‘The influence of aluminium on the ferrite formation and microstructural development in hot rolled dual-phase steel l’, Mater. Charact., 64 (2012) 69-78. 12. S. Shanmugam, R.D.K. Misra, T. Mannering, D. Panda, S.G. Jansto: ‘Impact toughness and microstructure relationship in niobium- and vanadium-microalloyed steels processed with varied cooling rates to similar yield strength’, Mater. Sci. Eng. A 437 (2006) 436-445. 13. Dae-Bum Park, Moo-Young Huh, Jae-Hyeok Shim, Jin-Yoo Suh, Kyu-Ho Lee, Woo-Sang Jung, Strengthening mechanism of hot rolled Ti and Nb microalloyed HSLA steels containing Mo and W with various coiling temperature, Mater. Sci. Eng. A, 2013; 560: 528-534. 14. Yong Woo Kim, Seok Weon Song, Seok Jong Seo, Seong-Gu Hong, Chong Soo Lee, Development of Ti and Mo micro-alloyed hot-rolled high strength sheet steel by controlling thermomechanical controlled processing schedule, Mater. Sci. Eng. A, 2013; 565: 430-438. 15. Y.-J. Zhang, G. Miyamoto, K. Shinbo, T. Furuhara, Effects of Į/Ȗ orientation relationship on VC interphase precipitation in low-carbon steels, Scripta Mater., 2013; 69: 17-20. 16. T, Gladman, Precipitation hardening in metals, Mater. Sci, Technol., 1999, 15: 30-36. 17. F. B. Pickering, Physical Metallurgy and the Design of Steels, Applied Science Publishing Ltd., London, 1978, 63. 18. .David A. Porter, and Kenneth E. Easterling, Phase Transformations in Metals and Alloys, 2rd ed. Chapman & Hall press, 1992.
Fig. 1 Optical micrographs and corresponding Vickers hardness distribution of Ti-Mo steel after different heat treatments: (a, e) isothermal holding at 625ºC for 5 min, (b, f) isothermal holding at 625ºC for 60 min, (c, g) isothermal holding at 725ºC for 5 min, (d, h) isothermal holding at 725ºC for 60 min. The number around each indentation represents the indentation test sequence. It is worthy to note that the size of the indentation does not change during the test sequence, which implies that the size of later indentations was not influenced by strain fields originating from previous indents. Fig. 2 TEM images of Ti-Mo bearing steel after isothermal holding at 625ºC for 60 min; (a) low- and (b, e, h) high-magnification TEM bright-field images showing two different precipitation morphologies (i.e., interface precipitation (zone A, zone B) and random homogeneous precipitation (zone C)) within the same polygonal ferrite grain, (c, f, i) high magnification dark field TEM images (using (200)TiMoC reflection of precipitation carbides, revealing the morphology of precipitates), and the (d, g, j) corresponding diffraction patterns, revealing the BN orientation relationship between precipitated carbides and ferrite matrix. Fig. 3 Schematic drawing illustrating the influence of isothermal holding temperature on the precipitation morphology within the ferrite matrix. Coexistence of different precipitation morphologies, namely interface precipitation and random homogeneous precipitation, can be observed in the HSLA steel that has undergone the general manufacturing process.
Figure1a-d
Figure1e-h
Figure2
Figure3
Table1
Table1 The microstructure characterization of Ti-Mo bearing steel after isothermal treatment at 625ºC and 725ºC for 5 and 60 min. Ti-Mo bearing steel presents all ferrite phase in the low isothermal holding temperature after very shortly isothermal holding times (5 min), indicating acceleration of the decomposition of austenite.
625кʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ʳ ˊ˅ˈк Isothermal holding time Ferrite Area fraction (%) Grain size (μm)
5 min 100 23.6̈́1.6
60 min 100 32.3̈́1.8
5 min
60 min
12.8̈́˄ˁ˅ʳ ʳ ʳ 28.7̈́˄ˁˉʳ ʳ .21.5̈́1.2
41.2̈́2.7
Table2
Table 2 Mechanical properties of Ti-Mo bearing steel after isothermal treatment at 625ºC for 5 and 60 min.
Process
Ys (MPa)
Ts (MPa)
A(%)
YR
625ʚ_5min
625ʚ_60min
Table3
Table 3 Components of yield stress in the steels studied.
Process
625
_5min
625 _60min
οߪ(MPa)
53.9
53.9
οߪ௦௦ (MPa)
.64.1
64.1
ο(MPa)
113.0
96.8
οߪௗ௦ (MPa)
168.0
168.0
οߪ௪ (MPa)
.274.0
243.2
ߪ௬ (MPa)
.673.0
626.0