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Effect of direct quenching on microstructure and mechanical properties of copper-bearing high-strength alloy steels
Materials Science and Engineering A252 (1998) 256 – 268
Effect of direct quenching on microstructure and mechanical properties of copper-bearing high-strength alloy steels Guen Chul Hwang a, Sunghak Lee *, Jang Yong Yoo b, Wung Yong Choo b a
b
Center for Ad6anced Aerospace Materials, Pohang Uni6ersity of Science and Technology, Pohang 790 -784, South Korea Plate, Rod and Welding Team, Technical Research Center, Pohang Iron and Steel Co, Ltd., Pohang, 790 -785, South Korea Received 11 March 1998
Abstract This study aims to investigate the effects of direct quenching on microstructural modification and mechanical properties of copper-bearing high-strength alloy steels. Two direct quenched and tempered (DQ&T) steel plates were rolled at different finish rolling temperatures, and their microstructures and mechanical properties were compared with those of a reheat quenched and tempered (RQ&T) steel plate. The as-quenched microstructure of the DQ plates consisted of refined lath martensite with high density of dislocations, which acted as preferred precipitation sites for NbC or o-Cu particles during tempering. These fine precipitates were not coarsened much up to the tempering temperature of about 650°C, and thus played a role in improving the tempering resistance. Especially in the DQ&T plate quenched at 760°C and tempered at 660°C, yield strength reached 1050 MPa, and Charpy impact energy at −18°C showed 140 J, indicating the potent effect of the DQ&T process. These findings indicated that the copper addition and the application of the DQ&T process to low-carbon alloy steels contributed to the production of steel plates with excellent strength and toughness. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Direct quenching; Reheat quenching; Microstructure
1. Introduction Hot rolling in the austenite temperature region has a grain refinement effect due to the repeated recrystallization, and produces rolled plates with more homogeneous compositions by gradually eliminating the segregation formed during casting. It enhances mechanical properties by breaking off or dispersing away the inevitable formation of non-metallic inclusions. Desired properties can also be optimally attained when the temperature and the rolling ratio in the controled rolling are precisely adjusted. Particularly in the case of low-carbon steels, controlled rolling has taken a more important role because both strength and toughness can be improved simultaneously by grain refinement. On-line direct quenching and tempering (DQ&T) process was recently developed as a part in the thermomechanical controled process (TMCP), and has been applied to high-strength steels. The DQ&T process has * Corresponding author. Tel.: + 82 562 2792715; fax: + 82 562 2792399. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(98)00670-4
several advantages over the conventional off-line reheat quenching and tempering (RQ&T) process: (1) the strength–toughness balance and the weldability can be enhanced because the microstructures and the precipitation behavior upon heat-treatment can be diversified, and (2) such heat-treatments as reheating and quenching can be skipped, thus lowering the manufacturing cost [1–4]. In a typical DQ&T process, the repeated recrystallization of austenite brings about the grain refinement by setting the finish rolling temperature at the austenite range, thereby making a fine quenched structure and introducing a considerable amount of dislocations in it. To make the quenched structure much finer, deformation bands are formed inside austenite grains before quenching by rolling in the two-phase region of austenite and ferrite [5]. The rolling parameters affecting the DQ&T microstructure are rolling temperature, rolling ratio, and cooling rate. By varying these parameters, a variety of microstructures can be achieved, and mechanical properties can be enhanced to a great extent by effectively making use of these varied microstruc-
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Fig. 1. Schematic diagrams of the processing conditions for (a) RQ&T and (b) DQ&T copper-bearing high-strength alloy steel plates.
tures. In the RQ&T process, on the contrary, the grain size and the internal structure are almost stabilized because the austenite structure is completely recrystallized during austenitization. Thus, the possibility to achieve microstructural modification by controlling the austenite structure is quite limited. Because hot rolling is in a non-equilibrium state thermodynamically, the phase solubility in austenite in the DQ&T process shows higher values than those under equilibrium at each temperature; hence, improving the quenching performance and increasing the amount of precipitates formed during subsequent tempering. However, the phase solubility in austenite approximates the values at equilibrium because of the lengthy austenitization at high temperature in the RQ&T process, thereby limiting the amount of precipitates. In this study, the DQ&T process was applied to a copper-bearing high-strength steel. In order to understand the mechanism of direct quenching, the correlation between microstructure and mechanical properties
of two DQ&T and one RQ&T conditions was investigated, and reported herein.
2. Experimental The steel used in this study is an HSLA-100 steel, a high-strength, high-toughness steel with a minimum yield strength of 690 MPa. The chemical composition is 0.04C–0.30Si–0.99Mn–0.005P–0.01S–3.71Ni–0.59Cr –1.80Cu–0.042Mo–0.048Al–0.039Nb. Although this is a low-carbon (0.04%C) alloy steel, its strength is significantly improved, together with other properties such as ductility, fracture toughness, and weldability, because fine o-Cu particles precipitate during tempering due to 1.8% Cu in its composition [6–11]. Copper atoms are sized by 0.2556 nm in diameter, showing only 3% difference compared with Fe atoms, and they are either dissolved in a-Fe or precipitated as o-Cu particles without forming intermetallic compounds with Fe. It is
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known that there is no interaction between Cu and interstitial C, O, and N [6]. It is also reported that the resistance to corrosion is enhanced by adding a small amount of copper to steel [7]. This copper-bearing high-strength steel was melted in a vacuum induction furnace and cast as a 27 kg ingot with a size of 120×120 ×250 mm. It was austenitized at 1200°C, and was forged to a 70-mm thick plate. After the forged plate was heated for 2 h at 1150 or 1250°C, the plate was rolled using a pilot plant rolling mill to a final thickness of about 13 mm. The rolling
Fig. 3. Optical micrographs of as-quenched (a) RQ, (b) DQ-760 and (c) DQ-855 plates, showing prior austenite grains.
Fig. 2. Optical micrographs of as-quenched (a) RQ, (b) DQ-760 and (c) DQ-855 plates. Nital etched.
ratio was in the range of 16–18% at each pass. The detailed DQ&T process is shown in Table 1. Rolling in the unrecrystallized austenite region was about 40– 60%, and brings the finish rolling temperature to 760°C in the case of austenitization at 1150°C and 855°C in the case at 1250°C. Immediately after rolling, the steel plate was quenched at a cooling rate of 8–14°C s − 1 in the DQ&T process. In comparison, in the RQ&T process, the plate was air cooled from 780°C to room temperature after rolling, austenitized again at 900°C for 1 h, and quenched at about 8°C s − 1. The specimens
Reheat temperature (°C)
1150
1250
Steel plate
DQ-760
DQ-855
1 2 3 4 5 6 7 8 9
1 2 3 4 5 6 7 8 9
Pass number
70.0 58.0 48.0 39.0 32.0 26.5 22.0 18.0 15.0
70.0 58.0 48.0 39.0 32.0 26.5 22.0 18.0 15.0 58.0 48.0 39.0 32.0 26.5 22.0 18.0 15.0 12.5
58.0 48.0 39.0 32.0 26.5 22.0 18.0 15 0 12.5 17.1 17.2 18.8 17.9 17.2 17.0 18.2 16.7 16.7
17.1 17.2 18.8 17.9 17.2 17.0 18.2 16.7 16.7
H1 (mm) H2 (mm) Reduction (%)
Table 1 Hot rolling schedules of direct quenched copper-bearing high-strength alloy steel plates
1205 1160 1105 1057 1005 978 950 905 855
1070 1040 1010 965 935 895 860 810 760 45.4 34.4 28.5 25.9 26.8 12.6 13.0 19.8 21.3
84.5 24.5 20.4 35.7 23.1 22.0 24.7 29.8 31.5 855
760
Temperature (°C) Inter-pass time (s) Finish rolling temperature (°C)
800
730
8.6
13.9
Start cooling tem- Cooling rate perature (°C) (°C s−1)
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were tempered for 1 h in the temperature range from 350 to 690°C. The two DQ processes, i.e. reheating at 1250°C and finish rolling at 855°C, and reheating at 1150°C and finish rolling at 760°C, are referred to as DQ-855 and DQ-760, respectively. After tempering, they are referred to as DQ-855&T and DQ-760&T, respectively. Fig. 1 is the schematic drawing of hot rolling and heat-treatment. Longitudinal round tensile specimens were sized by 6.25-mm in gage diameter and 25-mm in gage length. Tensile tests were conducted using a 10-t Instron machine at room temperature at a constant crosshead speed of 1 mm min − 1. Full-size Charpy impact specimens were machined with their notch perpendicular to the rolling direction (L – T direction), and impact tests were done at − 18 and −85°C. Fractured specimens were investigated by a scanning electron microscope (SEM). Size and shape of prior austenite grains were examined by an optical microscope after they were etched in Teepole solution. The tempered martensite structure including precipitates and carbides was observed by a transmission electron microscope (TEM) at an acceleration voltage of 300 kV. Carbon extraction replicas were also fabricated, and quantitative analyses of precipitates in various tempering conditions were conducted by a TEM.
3. Results
3.1. Microstructure Fig. 2(a)–(c) are optical micrographs of the reheat and quenched (RQ) and the direct quenched (DQ-760 and DQ-855) copper-bearing high-strength steel plates. All are composed of fine lath-type martensite and a small amount of inclusions like MnS as in other highstrength steels [12]. RQ plate has coarser and irregularly spaced laths (Fig. 2(a)) compared with DQ plates. Fig. 3(a)–(c) are optical micrographs showing prior austenite grains. Prior austenite grains of all the plates are sized by about 12 – 13 mm. In DQ plates, they are elongated in parallel to the rolling direction within which a large amount of deformation is present (Fig. 3(b) and (c)), whereas many equiaxed grains are observed in RQ plate (Fig. 3(a)). TEM micrographs of these as-quenched plates are shown in Fig. 4(a) – (c). Martensite is of lath-type, and the width of laths descends in the order of RQ, DQ855, and DQ-760 plates. A few fine NbC carbides are also observed, and their bright field image, dark field image, and EDS spectrum are shown in Fig. 5(a)–(c). When the plates are tempered, NbC carbides are newly precipitated or coarsened as the tempering temperature rises. TEM micrographs of carbon extraction replicas
Fig. 4. TEM micrographs of as-quenched (a) RQ, (b) DQ-760 and (c) DQ-855 plates, showing martensite laths.
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the plates tempered at 500°C, but no cementites are observed. In the case of DQ&T plates, precipitates are finer and more numerous than those in RQ&T plate. Between two DQ&T plates, DQ-760&T plate has finer precipitates than DQ-855&T plate. As the tempering temperature rises, NbC carbides get coarsened. In the case of RQ&T plate, they grow up to about 25 nm, whereas growth is less obvious in DQ&T plates due to the slower coarsening rate. Since it is almost impossible to extract o-Cu precipitates, the other fine particles in copper-bearing steels, in carbon extraction replicas, thin foil specimens were prepared for TEM observation. Fig. 7(a)–(f) are bright field and dark field TEM micrographs of DQ-760&T and RQ&T plates. In DQ-760&T plate tempered at 500°C, large amounts of very fine o-Cu precipitates and NbC carbides are distributed in a mixture. Because these two particles have similar shape and size, it is hard to distinguish them by appearance. Accordingly, they should be confirmed by dark field images obtained from microdiffraction patterns (Fig. 7(a) and (b)). A large number of dislocations are formed here, among which numerous fine particles are precipitated, thus suggesting that these dislocations provide precipitate nucleation sites. With increasing tempering temperature, it can be observed that o-Cu precipitates get coarsened and that dislocation density decreases (Fig. 7(c) and (d)). Precipitates in RQ&T plate are larger and less numerous than those in DQ-760&T plate (Fig. 7(e) and (f)). Table 2 lists the quantitative analysis data of such microstructural factors as prior austenite grain size, lath spacing, and size, volume fraction and spacing of precipitates. As the tempering temperature increases in both DQ&T and RQ&T plates, the size of fine NbC and o-Cu precipitates increases, but at a slower rate in DQ&T plate. At the same tempering temperature, the volume fraction of fine precipitates does not show much difference, but the number of particles tends to increase, together with the smaller particle spacing, in DQ&T plate because the prior austenite grains are smaller in DQ&T plate than in RQ&T plate. Particularly in DQ-760&T plate, this tendency is clear, and the Orowan effect due to the decreasing growth rate of fine particles greatly contributes to the strengthening of copper-bearing high-strength steels.
3.2. Mechanical properties
Fig. 5. (a) Bright field image, (b) dark field image and (c) EDS spectra of NbC precipitates in as-quenched RQ plate.
of these three plates tempered at 500, 600 and 660°C, respectively, are shown in Fig. 6(a) – (i). Extremely fine NbC precipitates sized less than 10 nm are observed in
Fig. 8 shows the hardness test results as a function of tempering temperature. In all the plates, hardness increases with increasing tempering temperature until it reaches the maximum at 500°C. After the maximum, hardness is reduced drastically, and shows a typical dependence of precipitates on temperature. Two DQ&T plates have similar hardnesses, higher than RQ&T plate’s.
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Fig. 6. TEM micrographs of carbon extraction replicas of (a) through (c) RQ&T, (d) through (f) DQ-760&T and (g) through (i) DQ-855&T plates, showing NbC precipitates. Each specimen is tempered at 500, 600 and 660°C, respectively, as marked in micrographs.
Fig. 9(a) and (b) are the tensile test results versus tempering temperature. Tensile strength reaches the maximum at 500°C as on the hardness test. For a given tempering temperature, it descends in the order of DQ-760&T, DQ-855&T and RQ&T plates (Fig. 9(a)). Yield strength also descends in the same order, but DQ-760&T and DQ-855&T plates maintain their strength up to quite high tempering temperatures (Fig. 9(b)). Elongation is maintained nearly constant, not being affected by tempering temperature, and shows higher values in RQ&T plate than in two DQ&T plates. It can be learned from these tensile results that DQ760&T plate has excellent tensile properties overall, holding up its yield strength at about 1050 MPa up to 660°C and showing high resistance to tempering with about 20% elongation. Fig. 10(a) and (b) illustrate the impact test results at
−18 and − 85°C versus tempering temperature. The impact tests were conducted at tempering temperatures over 500°C, higher than the 300–400°C tempered martensite embrittlement (TME) region. As the tempering temperature rises, impact energy values increase, with lower impact energy values measured at −85°C than that measured at − 18°C. Impact energy of RQ&T plate tested at −85°C abruptly increases with increasing the tempering temperature. When tempered at 660°C, impact energy reaches a similar level to that at −18°C, indicating that the ductile-brittle transition temperature is below − 85°C. Impact energy ascends in the order of DQ-760&T, DQ-855&T, and RQ&T plates, contrary to the tendency in tensile and yield strengths. This indicates that impact energy is reduced by the increased portion in strengths due to the DQ&T process.
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Fig. 7. TEM micrographs of (a) DQ-760&T plate tempered at 500°C, (c) DQ-760&T plate tempered at 660°C, and (e) RQ&T plate tempered at 660°C. (b), (d), and (f) are dark field images of (a), (c), and (e), respectively, obtained from the microdiffraction patterns of o-Cu precipitates.
3.3. Fracture surface
4. Discussion
Fig. 11(a) and (b) are fractographs of tensile specimens of RQ&T and DQ-760&T plates tempered at 500°C. The major fracture mode is ductile fracture consisted of fine dimples. The dimple size of asquenched and as-tempered plates was measured, and lists in Table 3. Dimples are the largest in as-quenched plates, the smallest when tempered at 500°C, and increased with increasing the tempering temperature. At the same tempering temperature, the dimple size of RQ&T plate is larger than that of DQ-760&T plate. This indicates that the precipitation of fine particles is related with ductile fracture, and that fine precipitates get smallest at the tempering temperature of 500°C at which hardness and strength show the maximum values. Dimples get larger as precipitates become gradually coarsened over 500°C.
In copper-bearing high-strength steels, dislocation strengthening due to lath martensite and precipitation strengthening due to NbC and o-Cu particles precipitated during tempering are major strengthening mechanisms. It is important to understand precipitation of fine particles, their growth, and correlation with dislocations. Hardness of typical martensites decreases almost linearly with increasing the tempering temperature [13]. Comparing with RQ&T plate, DQ&T plates requires higher driving force for recrystallization or for recovery due to many defects formed during hot rolling. Thus, the matrix is supposed to get softened much faster, but this matrix softening is delayed by the significant effect of sub-boundary pinning [14] due to fine precipitates. As the tempering temperature rises, fine precipitates grow slower in DQ&T plate than in
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Fig. 7. (Continued)
RQ&T plate, and the precipitate spacing is denser, resulting in high strength and hardness over the whole tempering temperature range. This Orowan-type strengthening is closely related to the change in strength and the dimple size on the fracture surface. The precipitation and growth behavior of o-Cu and NbC particles vary with the fabrication process of DQ&T and RQ&T plates. In DQ&T plate, deformation state of austenite endowed in hot rolling before quenching persists down until room temperature, raising the dislocation density of room-temperature structure. As a result, more nucleation sites are available for more fine particles to precipitate during tempering (Fig. 7(a)). Because fine precipitates grow by diffusion, they get smaller when the diffusion radius of copper atoms is reduced, and thus finer precipitates are formed at the initial precipitation stage in DQ&T plate than in RQ&T plate. The quantitative data in Table 2 show that precipitates of DQ&T plates are much smaller than those of RQ&T plate at the tempering temperature of 660°C and that they rarely
grow compared with the precipitate size at 500°C. This phenomenon results from the slow growth rate of precipitates in DQ&T plate, and can be attributed to the decreasing driving force for the precipitate growth because of reducing interfacial energy between precipitates and the matrix when precipitates get smaller. Accordingly, fine precipitates in DQ&T plate grow slowly and effectively prevent the dislocation movement, thereby improving the tempering resistance and maintaining the maximum yield strength up to around 650°C (Fig. 9(b)). The high yield strength over such a wide tempering temperature range is because particles precipitated fine inside the matrix do not get any coarser even if dislocations are abruptly extinguished. Fig. 12(a) and (b) illustrate the correlation of yield strength and impact energy to compare overall mechanical properties. DQ&T plate hasmore excellent properties than RQ&T plate, irrespective of the test temperature. Particularly, DQ760&T plate has the most excellent mechanical properties because of higher strength and relatively smaller reduction of impact energy than DQ-855&T plate.
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Table 2 Quantitative metallographic results Steel plate
Tempering temperature (°C)
Prior g grain size Lath spacing (mm) (mm)
Volume fraction of NbCa (%)
Particle sizea (nm)
NbCb
o-Cuc
Spacing of NbCb (nm)
RQ&T
— 500 600 660
12.5 — — —
0.33 — — —
1.8 1.9 2.1 2.1
10 12 17 25
— 5 13 27
61 64 76 98
DQ-760&T
— 500 600 660
12.0 — —
0.19 — — —
— 1.3 1.6 1.7
— 3 4 5
— 4 6 8
— 23 25 28
DQ-855&T
— 500 600 660
13.0 — — —
0.24 — — —
— 1.4 1.7 1.9
— 4 5 6
— 5 7 10
— 27 37 40
a
Measured values might be overestimated because of the vagueness of the micrographs. Quantitatively measured from TEM micrographs of carbon extraction replicas. c Quantitatively measured from TEM micrographs of thin foil specimens. b
By adding copper to low-carbon alloy steels and processing them in the DQ&T process, high-strength steels with excellent strength and toughness can be produced. Especially, DQ&T steel plates satisfactorily meet with the specifications of HSLA-100 steel, and have excellent weldability due to the low carbon content of 0.04%. Consequently, the results of the present study on the fabrication of low-carbon high-strength alloy steel plates with fine grains and laths and with excellent tempering resistance of fine precipitates by the DQ&T process will help better understand the refinement mechanism occurring during hot rolling. It is
Fig. 8. Vickers hardness vs tempering temperature for RQ&T, DQ760&T and DQ-855&T plates.
further expected that they will be used as important experimental data to develop high-strength, high-toughness steels by adding alloying elements and by applying controlled hot rolling process to low-carbon alloy steels.
5. Conclusions In the present study, the DQ&T process was applied to low-carbon copper-bearing alloy steels, and the resultant correlation between microstructure and mechanical properties was investigated. The conclusions are summarized as follows: 1. In the as-quenched DQ plates, prior austenite grains are elongated parallel to the rolling direction, within which many deformed structures are present, whereas the as-quenched RQ plate mainly consists of equiaxed grains. When these plates are tempered, fine NbC and o-Cu particles are precipitated with increasing the tempering temperature, and DQ&T plates have finer precipitates than RQ&T plate. 2. Tensile and yield strengths are the highest at the tempering temperature of 500°C. At a given tempering temperature, they decrease in the order of DQ-760&T, DQ-855&T, and RQ&T plates. Yield strength of DQ&T plates is maintained up to 660°C, showing an excellent tempering resistance. Ductility and impact energy increase in the order of DQ-760&T, DQ-855&T and RQ&T plates. 3. The structure of the low-carbon copper-bearing alloy steel plate fabricated by the DQ&T process has been transformed to a fine one with high dislocation
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Fig. 9. (a) Tensile strength and (b) yield strength and elongation vs tempering temperature for RQ&T, DQ-760&T and DQ-855&T plates.
density. During tempering, fine NbC and o-Cu particles are precipitated in large amounts, which do not get coarsened even when the tempering temperature rises, resulting in excellent mechanical properties. Results of
this study indicate that the addition of alloying elements and the application of the DQ&T process to low-carbon alloy steel plates contribute to the production of plates with excellent strength and toughness.
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Acknowledgements This work has been supported by the Agency for Defense Development. The authors would like to thank Young Woo Kim and Dr In Ok Shim of Agency for Defense Development, and Yong Gyu Choi and Hyang Jin Koh of POSTECH for their helpful discussion on microstructural analysis.
Fig. 11. SEM fractographs of the fractured tensile specimens of (a) RQ&T and (b) DQ-760&T plates tempered at 500°C, showing ductile dimpled fracture mode.
Table 3 Dimple size measured from the fractographic results
Fig. 10. Charpy impact energy at (a)−18 and (b)− 85°C vs tempering temperature for RQ&T, DQ-760&T and DQ-0855&T plates.
Steel plate
Tempering temperature (°C)
Dimple size (mm)
RQ&T
— 500 600 660 690
2.9 1.6 2.1 2.2 2.3
DQ-760&T
— 500 600 660 690
2.1 1.0 1.8 1.9 2.0
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References [1] J. Prohaszka, J. Dobranszky, J. Heat Treating 9 (1991) 63. [2] M.T. Miglin, J.P. Hirth, A.R. Rosenfield, W.A.T. Clark, Metall. Trans. A 17A (1986) 791. [3] G.E. Hicho, R.J. Fields, J. Heat Treating 8 (1990) 101. [4] G. Ronchato, M. Castagna, R.L. Colombo, J. Heat Treating 4 (1985) 194. [5] S. Sangal, S. Yannacopoulos, Can. Metall. Q. 31 (1992) 55. [6] E.C. Bain, H.W. Paxton, Alloying Elements in Steel, 2nd edn, ASM, Metals Park, OH, 1962, p. 59. [7] J.C. West, J. Ship. Prod. 3 (1987) 111. [8] B. Dutta, C.M. Sellars, Mater. Sci. Technol. 3 (1987) 197. [9] T. Abe, M. Kurihara, H. Tagawa, K. Tsukada, Trans. Iron Steel Inst. Jpn. 27 (1987) 478. [10] E. Hornbogen, R.C. Glenn, Trans. AIME 218 (1960) 1064. [11] P.G. Shewmon, Transformation in Metals, McGraw-Hill, New York, NY, 1969, p. 299. [12] J.R. Wilcox, R.W.K. Honeycombe, Mater. Sci. Technol. 3 (1987) 849. [13] G.R. Speich, T.M. Scoonover, in: A.J. DeArdo (Ed.), Processing, Microstructure, and Properties of HSLA Steels, TMS, Warrendale, PA, 1988, p. 263. [14] J.W. Martin, Micromechanisms in Particle-Hardened Alloys, Cambridge University Press, New York, NY, 1980, p. 151.
Fig. 12. Charpy impact energy at (a)− 18 and (b)− 85°C vs yield strength for RQ&T, DQ-760&T and DQ-855&T plates tempered at 500, 600, 630 and 660°C.
.
Effect of direct quenching on microstructure and mechanical properties of copper-bearing high-strength alloy steels Guen Chul Hwang a, Sunghak Lee *, Jang Yong Yoo b, Wung Yong Choo b a
b
Center for Ad6anced Aerospace Materials, Pohang Uni6ersity of Science and Technology, Pohang 790 -784, South Korea Plate, Rod and Welding Team, Technical Research Center, Pohang Iron and Steel Co, Ltd., Pohang, 790 -785, South Korea Received 11 March 1998
Abstract This study aims to investigate the effects of direct quenching on microstructural modification and mechanical properties of copper-bearing high-strength alloy steels. Two direct quenched and tempered (DQ&T) steel plates were rolled at different finish rolling temperatures, and their microstructures and mechanical properties were compared with those of a reheat quenched and tempered (RQ&T) steel plate. The as-quenched microstructure of the DQ plates consisted of refined lath martensite with high density of dislocations, which acted as preferred precipitation sites for NbC or o-Cu particles during tempering. These fine precipitates were not coarsened much up to the tempering temperature of about 650°C, and thus played a role in improving the tempering resistance. Especially in the DQ&T plate quenched at 760°C and tempered at 660°C, yield strength reached 1050 MPa, and Charpy impact energy at −18°C showed 140 J, indicating the potent effect of the DQ&T process. These findings indicated that the copper addition and the application of the DQ&T process to low-carbon alloy steels contributed to the production of steel plates with excellent strength and toughness. © 1998 Elsevier Science S.A. All rights reserved. Keywords: Direct quenching; Reheat quenching; Microstructure
1. Introduction Hot rolling in the austenite temperature region has a grain refinement effect due to the repeated recrystallization, and produces rolled plates with more homogeneous compositions by gradually eliminating the segregation formed during casting. It enhances mechanical properties by breaking off or dispersing away the inevitable formation of non-metallic inclusions. Desired properties can also be optimally attained when the temperature and the rolling ratio in the controled rolling are precisely adjusted. Particularly in the case of low-carbon steels, controlled rolling has taken a more important role because both strength and toughness can be improved simultaneously by grain refinement. On-line direct quenching and tempering (DQ&T) process was recently developed as a part in the thermomechanical controled process (TMCP), and has been applied to high-strength steels. The DQ&T process has * Corresponding author. Tel.: + 82 562 2792715; fax: + 82 562 2792399. 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(98)00670-4
several advantages over the conventional off-line reheat quenching and tempering (RQ&T) process: (1) the strength–toughness balance and the weldability can be enhanced because the microstructures and the precipitation behavior upon heat-treatment can be diversified, and (2) such heat-treatments as reheating and quenching can be skipped, thus lowering the manufacturing cost [1–4]. In a typical DQ&T process, the repeated recrystallization of austenite brings about the grain refinement by setting the finish rolling temperature at the austenite range, thereby making a fine quenched structure and introducing a considerable amount of dislocations in it. To make the quenched structure much finer, deformation bands are formed inside austenite grains before quenching by rolling in the two-phase region of austenite and ferrite [5]. The rolling parameters affecting the DQ&T microstructure are rolling temperature, rolling ratio, and cooling rate. By varying these parameters, a variety of microstructures can be achieved, and mechanical properties can be enhanced to a great extent by effectively making use of these varied microstruc-
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Fig. 1. Schematic diagrams of the processing conditions for (a) RQ&T and (b) DQ&T copper-bearing high-strength alloy steel plates.
tures. In the RQ&T process, on the contrary, the grain size and the internal structure are almost stabilized because the austenite structure is completely recrystallized during austenitization. Thus, the possibility to achieve microstructural modification by controlling the austenite structure is quite limited. Because hot rolling is in a non-equilibrium state thermodynamically, the phase solubility in austenite in the DQ&T process shows higher values than those under equilibrium at each temperature; hence, improving the quenching performance and increasing the amount of precipitates formed during subsequent tempering. However, the phase solubility in austenite approximates the values at equilibrium because of the lengthy austenitization at high temperature in the RQ&T process, thereby limiting the amount of precipitates. In this study, the DQ&T process was applied to a copper-bearing high-strength steel. In order to understand the mechanism of direct quenching, the correlation between microstructure and mechanical properties
of two DQ&T and one RQ&T conditions was investigated, and reported herein.
2. Experimental The steel used in this study is an HSLA-100 steel, a high-strength, high-toughness steel with a minimum yield strength of 690 MPa. The chemical composition is 0.04C–0.30Si–0.99Mn–0.005P–0.01S–3.71Ni–0.59Cr –1.80Cu–0.042Mo–0.048Al–0.039Nb. Although this is a low-carbon (0.04%C) alloy steel, its strength is significantly improved, together with other properties such as ductility, fracture toughness, and weldability, because fine o-Cu particles precipitate during tempering due to 1.8% Cu in its composition [6–11]. Copper atoms are sized by 0.2556 nm in diameter, showing only 3% difference compared with Fe atoms, and they are either dissolved in a-Fe or precipitated as o-Cu particles without forming intermetallic compounds with Fe. It is
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known that there is no interaction between Cu and interstitial C, O, and N [6]. It is also reported that the resistance to corrosion is enhanced by adding a small amount of copper to steel [7]. This copper-bearing high-strength steel was melted in a vacuum induction furnace and cast as a 27 kg ingot with a size of 120×120 ×250 mm. It was austenitized at 1200°C, and was forged to a 70-mm thick plate. After the forged plate was heated for 2 h at 1150 or 1250°C, the plate was rolled using a pilot plant rolling mill to a final thickness of about 13 mm. The rolling
Fig. 3. Optical micrographs of as-quenched (a) RQ, (b) DQ-760 and (c) DQ-855 plates, showing prior austenite grains.
Fig. 2. Optical micrographs of as-quenched (a) RQ, (b) DQ-760 and (c) DQ-855 plates. Nital etched.
ratio was in the range of 16–18% at each pass. The detailed DQ&T process is shown in Table 1. Rolling in the unrecrystallized austenite region was about 40– 60%, and brings the finish rolling temperature to 760°C in the case of austenitization at 1150°C and 855°C in the case at 1250°C. Immediately after rolling, the steel plate was quenched at a cooling rate of 8–14°C s − 1 in the DQ&T process. In comparison, in the RQ&T process, the plate was air cooled from 780°C to room temperature after rolling, austenitized again at 900°C for 1 h, and quenched at about 8°C s − 1. The specimens
Reheat temperature (°C)
1150
1250
Steel plate
DQ-760
DQ-855
1 2 3 4 5 6 7 8 9
1 2 3 4 5 6 7 8 9
Pass number
70.0 58.0 48.0 39.0 32.0 26.5 22.0 18.0 15.0
70.0 58.0 48.0 39.0 32.0 26.5 22.0 18.0 15.0 58.0 48.0 39.0 32.0 26.5 22.0 18.0 15.0 12.5
58.0 48.0 39.0 32.0 26.5 22.0 18.0 15 0 12.5 17.1 17.2 18.8 17.9 17.2 17.0 18.2 16.7 16.7
17.1 17.2 18.8 17.9 17.2 17.0 18.2 16.7 16.7
H1 (mm) H2 (mm) Reduction (%)
Table 1 Hot rolling schedules of direct quenched copper-bearing high-strength alloy steel plates
1205 1160 1105 1057 1005 978 950 905 855
1070 1040 1010 965 935 895 860 810 760 45.4 34.4 28.5 25.9 26.8 12.6 13.0 19.8 21.3
84.5 24.5 20.4 35.7 23.1 22.0 24.7 29.8 31.5 855
760
Temperature (°C) Inter-pass time (s) Finish rolling temperature (°C)
800
730
8.6
13.9
Start cooling tem- Cooling rate perature (°C) (°C s−1)
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were tempered for 1 h in the temperature range from 350 to 690°C. The two DQ processes, i.e. reheating at 1250°C and finish rolling at 855°C, and reheating at 1150°C and finish rolling at 760°C, are referred to as DQ-855 and DQ-760, respectively. After tempering, they are referred to as DQ-855&T and DQ-760&T, respectively. Fig. 1 is the schematic drawing of hot rolling and heat-treatment. Longitudinal round tensile specimens were sized by 6.25-mm in gage diameter and 25-mm in gage length. Tensile tests were conducted using a 10-t Instron machine at room temperature at a constant crosshead speed of 1 mm min − 1. Full-size Charpy impact specimens were machined with their notch perpendicular to the rolling direction (L – T direction), and impact tests were done at − 18 and −85°C. Fractured specimens were investigated by a scanning electron microscope (SEM). Size and shape of prior austenite grains were examined by an optical microscope after they were etched in Teepole solution. The tempered martensite structure including precipitates and carbides was observed by a transmission electron microscope (TEM) at an acceleration voltage of 300 kV. Carbon extraction replicas were also fabricated, and quantitative analyses of precipitates in various tempering conditions were conducted by a TEM.
3. Results
3.1. Microstructure Fig. 2(a)–(c) are optical micrographs of the reheat and quenched (RQ) and the direct quenched (DQ-760 and DQ-855) copper-bearing high-strength steel plates. All are composed of fine lath-type martensite and a small amount of inclusions like MnS as in other highstrength steels [12]. RQ plate has coarser and irregularly spaced laths (Fig. 2(a)) compared with DQ plates. Fig. 3(a)–(c) are optical micrographs showing prior austenite grains. Prior austenite grains of all the plates are sized by about 12 – 13 mm. In DQ plates, they are elongated in parallel to the rolling direction within which a large amount of deformation is present (Fig. 3(b) and (c)), whereas many equiaxed grains are observed in RQ plate (Fig. 3(a)). TEM micrographs of these as-quenched plates are shown in Fig. 4(a) – (c). Martensite is of lath-type, and the width of laths descends in the order of RQ, DQ855, and DQ-760 plates. A few fine NbC carbides are also observed, and their bright field image, dark field image, and EDS spectrum are shown in Fig. 5(a)–(c). When the plates are tempered, NbC carbides are newly precipitated or coarsened as the tempering temperature rises. TEM micrographs of carbon extraction replicas
Fig. 4. TEM micrographs of as-quenched (a) RQ, (b) DQ-760 and (c) DQ-855 plates, showing martensite laths.
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the plates tempered at 500°C, but no cementites are observed. In the case of DQ&T plates, precipitates are finer and more numerous than those in RQ&T plate. Between two DQ&T plates, DQ-760&T plate has finer precipitates than DQ-855&T plate. As the tempering temperature rises, NbC carbides get coarsened. In the case of RQ&T plate, they grow up to about 25 nm, whereas growth is less obvious in DQ&T plates due to the slower coarsening rate. Since it is almost impossible to extract o-Cu precipitates, the other fine particles in copper-bearing steels, in carbon extraction replicas, thin foil specimens were prepared for TEM observation. Fig. 7(a)–(f) are bright field and dark field TEM micrographs of DQ-760&T and RQ&T plates. In DQ-760&T plate tempered at 500°C, large amounts of very fine o-Cu precipitates and NbC carbides are distributed in a mixture. Because these two particles have similar shape and size, it is hard to distinguish them by appearance. Accordingly, they should be confirmed by dark field images obtained from microdiffraction patterns (Fig. 7(a) and (b)). A large number of dislocations are formed here, among which numerous fine particles are precipitated, thus suggesting that these dislocations provide precipitate nucleation sites. With increasing tempering temperature, it can be observed that o-Cu precipitates get coarsened and that dislocation density decreases (Fig. 7(c) and (d)). Precipitates in RQ&T plate are larger and less numerous than those in DQ-760&T plate (Fig. 7(e) and (f)). Table 2 lists the quantitative analysis data of such microstructural factors as prior austenite grain size, lath spacing, and size, volume fraction and spacing of precipitates. As the tempering temperature increases in both DQ&T and RQ&T plates, the size of fine NbC and o-Cu precipitates increases, but at a slower rate in DQ&T plate. At the same tempering temperature, the volume fraction of fine precipitates does not show much difference, but the number of particles tends to increase, together with the smaller particle spacing, in DQ&T plate because the prior austenite grains are smaller in DQ&T plate than in RQ&T plate. Particularly in DQ-760&T plate, this tendency is clear, and the Orowan effect due to the decreasing growth rate of fine particles greatly contributes to the strengthening of copper-bearing high-strength steels.
3.2. Mechanical properties
Fig. 5. (a) Bright field image, (b) dark field image and (c) EDS spectra of NbC precipitates in as-quenched RQ plate.
of these three plates tempered at 500, 600 and 660°C, respectively, are shown in Fig. 6(a) – (i). Extremely fine NbC precipitates sized less than 10 nm are observed in
Fig. 8 shows the hardness test results as a function of tempering temperature. In all the plates, hardness increases with increasing tempering temperature until it reaches the maximum at 500°C. After the maximum, hardness is reduced drastically, and shows a typical dependence of precipitates on temperature. Two DQ&T plates have similar hardnesses, higher than RQ&T plate’s.
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Fig. 6. TEM micrographs of carbon extraction replicas of (a) through (c) RQ&T, (d) through (f) DQ-760&T and (g) through (i) DQ-855&T plates, showing NbC precipitates. Each specimen is tempered at 500, 600 and 660°C, respectively, as marked in micrographs.
Fig. 9(a) and (b) are the tensile test results versus tempering temperature. Tensile strength reaches the maximum at 500°C as on the hardness test. For a given tempering temperature, it descends in the order of DQ-760&T, DQ-855&T and RQ&T plates (Fig. 9(a)). Yield strength also descends in the same order, but DQ-760&T and DQ-855&T plates maintain their strength up to quite high tempering temperatures (Fig. 9(b)). Elongation is maintained nearly constant, not being affected by tempering temperature, and shows higher values in RQ&T plate than in two DQ&T plates. It can be learned from these tensile results that DQ760&T plate has excellent tensile properties overall, holding up its yield strength at about 1050 MPa up to 660°C and showing high resistance to tempering with about 20% elongation. Fig. 10(a) and (b) illustrate the impact test results at
−18 and − 85°C versus tempering temperature. The impact tests were conducted at tempering temperatures over 500°C, higher than the 300–400°C tempered martensite embrittlement (TME) region. As the tempering temperature rises, impact energy values increase, with lower impact energy values measured at −85°C than that measured at − 18°C. Impact energy of RQ&T plate tested at −85°C abruptly increases with increasing the tempering temperature. When tempered at 660°C, impact energy reaches a similar level to that at −18°C, indicating that the ductile-brittle transition temperature is below − 85°C. Impact energy ascends in the order of DQ-760&T, DQ-855&T, and RQ&T plates, contrary to the tendency in tensile and yield strengths. This indicates that impact energy is reduced by the increased portion in strengths due to the DQ&T process.
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Fig. 7. TEM micrographs of (a) DQ-760&T plate tempered at 500°C, (c) DQ-760&T plate tempered at 660°C, and (e) RQ&T plate tempered at 660°C. (b), (d), and (f) are dark field images of (a), (c), and (e), respectively, obtained from the microdiffraction patterns of o-Cu precipitates.
3.3. Fracture surface
4. Discussion
Fig. 11(a) and (b) are fractographs of tensile specimens of RQ&T and DQ-760&T plates tempered at 500°C. The major fracture mode is ductile fracture consisted of fine dimples. The dimple size of asquenched and as-tempered plates was measured, and lists in Table 3. Dimples are the largest in as-quenched plates, the smallest when tempered at 500°C, and increased with increasing the tempering temperature. At the same tempering temperature, the dimple size of RQ&T plate is larger than that of DQ-760&T plate. This indicates that the precipitation of fine particles is related with ductile fracture, and that fine precipitates get smallest at the tempering temperature of 500°C at which hardness and strength show the maximum values. Dimples get larger as precipitates become gradually coarsened over 500°C.
In copper-bearing high-strength steels, dislocation strengthening due to lath martensite and precipitation strengthening due to NbC and o-Cu particles precipitated during tempering are major strengthening mechanisms. It is important to understand precipitation of fine particles, their growth, and correlation with dislocations. Hardness of typical martensites decreases almost linearly with increasing the tempering temperature [13]. Comparing with RQ&T plate, DQ&T plates requires higher driving force for recrystallization or for recovery due to many defects formed during hot rolling. Thus, the matrix is supposed to get softened much faster, but this matrix softening is delayed by the significant effect of sub-boundary pinning [14] due to fine precipitates. As the tempering temperature rises, fine precipitates grow slower in DQ&T plate than in
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Fig. 7. (Continued)
RQ&T plate, and the precipitate spacing is denser, resulting in high strength and hardness over the whole tempering temperature range. This Orowan-type strengthening is closely related to the change in strength and the dimple size on the fracture surface. The precipitation and growth behavior of o-Cu and NbC particles vary with the fabrication process of DQ&T and RQ&T plates. In DQ&T plate, deformation state of austenite endowed in hot rolling before quenching persists down until room temperature, raising the dislocation density of room-temperature structure. As a result, more nucleation sites are available for more fine particles to precipitate during tempering (Fig. 7(a)). Because fine precipitates grow by diffusion, they get smaller when the diffusion radius of copper atoms is reduced, and thus finer precipitates are formed at the initial precipitation stage in DQ&T plate than in RQ&T plate. The quantitative data in Table 2 show that precipitates of DQ&T plates are much smaller than those of RQ&T plate at the tempering temperature of 660°C and that they rarely
grow compared with the precipitate size at 500°C. This phenomenon results from the slow growth rate of precipitates in DQ&T plate, and can be attributed to the decreasing driving force for the precipitate growth because of reducing interfacial energy between precipitates and the matrix when precipitates get smaller. Accordingly, fine precipitates in DQ&T plate grow slowly and effectively prevent the dislocation movement, thereby improving the tempering resistance and maintaining the maximum yield strength up to around 650°C (Fig. 9(b)). The high yield strength over such a wide tempering temperature range is because particles precipitated fine inside the matrix do not get any coarser even if dislocations are abruptly extinguished. Fig. 12(a) and (b) illustrate the correlation of yield strength and impact energy to compare overall mechanical properties. DQ&T plate hasmore excellent properties than RQ&T plate, irrespective of the test temperature. Particularly, DQ760&T plate has the most excellent mechanical properties because of higher strength and relatively smaller reduction of impact energy than DQ-855&T plate.
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Table 2 Quantitative metallographic results Steel plate
Tempering temperature (°C)
Prior g grain size Lath spacing (mm) (mm)
Volume fraction of NbCa (%)
Particle sizea (nm)
NbCb
o-Cuc
Spacing of NbCb (nm)
RQ&T
— 500 600 660
12.5 — — —
0.33 — — —
1.8 1.9 2.1 2.1
10 12 17 25
— 5 13 27
61 64 76 98
DQ-760&T
— 500 600 660
12.0 — —
0.19 — — —
— 1.3 1.6 1.7
— 3 4 5
— 4 6 8
— 23 25 28
DQ-855&T
— 500 600 660
13.0 — — —
0.24 — — —
— 1.4 1.7 1.9
— 4 5 6
— 5 7 10
— 27 37 40
a
Measured values might be overestimated because of the vagueness of the micrographs. Quantitatively measured from TEM micrographs of carbon extraction replicas. c Quantitatively measured from TEM micrographs of thin foil specimens. b
By adding copper to low-carbon alloy steels and processing them in the DQ&T process, high-strength steels with excellent strength and toughness can be produced. Especially, DQ&T steel plates satisfactorily meet with the specifications of HSLA-100 steel, and have excellent weldability due to the low carbon content of 0.04%. Consequently, the results of the present study on the fabrication of low-carbon high-strength alloy steel plates with fine grains and laths and with excellent tempering resistance of fine precipitates by the DQ&T process will help better understand the refinement mechanism occurring during hot rolling. It is
Fig. 8. Vickers hardness vs tempering temperature for RQ&T, DQ760&T and DQ-855&T plates.
further expected that they will be used as important experimental data to develop high-strength, high-toughness steels by adding alloying elements and by applying controlled hot rolling process to low-carbon alloy steels.
5. Conclusions In the present study, the DQ&T process was applied to low-carbon copper-bearing alloy steels, and the resultant correlation between microstructure and mechanical properties was investigated. The conclusions are summarized as follows: 1. In the as-quenched DQ plates, prior austenite grains are elongated parallel to the rolling direction, within which many deformed structures are present, whereas the as-quenched RQ plate mainly consists of equiaxed grains. When these plates are tempered, fine NbC and o-Cu particles are precipitated with increasing the tempering temperature, and DQ&T plates have finer precipitates than RQ&T plate. 2. Tensile and yield strengths are the highest at the tempering temperature of 500°C. At a given tempering temperature, they decrease in the order of DQ-760&T, DQ-855&T, and RQ&T plates. Yield strength of DQ&T plates is maintained up to 660°C, showing an excellent tempering resistance. Ductility and impact energy increase in the order of DQ-760&T, DQ-855&T and RQ&T plates. 3. The structure of the low-carbon copper-bearing alloy steel plate fabricated by the DQ&T process has been transformed to a fine one with high dislocation
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Fig. 9. (a) Tensile strength and (b) yield strength and elongation vs tempering temperature for RQ&T, DQ-760&T and DQ-855&T plates.
density. During tempering, fine NbC and o-Cu particles are precipitated in large amounts, which do not get coarsened even when the tempering temperature rises, resulting in excellent mechanical properties. Results of
this study indicate that the addition of alloying elements and the application of the DQ&T process to low-carbon alloy steel plates contribute to the production of plates with excellent strength and toughness.
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Acknowledgements This work has been supported by the Agency for Defense Development. The authors would like to thank Young Woo Kim and Dr In Ok Shim of Agency for Defense Development, and Yong Gyu Choi and Hyang Jin Koh of POSTECH for their helpful discussion on microstructural analysis.
Fig. 11. SEM fractographs of the fractured tensile specimens of (a) RQ&T and (b) DQ-760&T plates tempered at 500°C, showing ductile dimpled fracture mode.
Table 3 Dimple size measured from the fractographic results
Fig. 10. Charpy impact energy at (a)−18 and (b)− 85°C vs tempering temperature for RQ&T, DQ-760&T and DQ-0855&T plates.
Steel plate
Tempering temperature (°C)
Dimple size (mm)
RQ&T
— 500 600 660 690
2.9 1.6 2.1 2.2 2.3
DQ-760&T
— 500 600 660 690
2.1 1.0 1.8 1.9 2.0
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Fig. 12. Charpy impact energy at (a)− 18 and (b)− 85°C vs yield strength for RQ&T, DQ-760&T and DQ-855&T plates tempered at 500, 600, 630 and 660°C.
.