2016 Volume 56 Issue 12 Pages 2284-2289
Study on continuous cooling transformation (CCT) behavior is an essential issue before thermo-mechanical processing for a new steel. In this study, dynamic CCT characteristics and microstructural evolution of a novel Cu-bearing pipeline steel with different Cu content (1.06%, 1.46% and 2.00%) were investigated by means of a combined method of dilatometry and metallography. The microstructure developed at a cooling rate range of 0.05 to 30°C/s consisted of pearlite, polygonal ferrite, quasi-polygonal ferrite and acicular ferrite. More Cu addition could lower the transformation temperature for austenite to ferrite and lead to an increase in the driving force for the acicular ferrite transformation, resulting in a full acicular ferrite for 2.0 Cu steel at cooling rate above 2°C/s. The precipitation behavior of Cu-rich phase during continuous cooling showed that Cu precipitation could occur in the acicular ferrite, which made a hardness peak on the hardness vs cooling rate curve of 2.0Cu steel at the cooling rate of 2°C/s. However, no Cu precipitate was detected in the acicular ferrite at higher cooling rate for 2.0 Cu steel. Higher supersaturation of Cu in austenite and a short incubation period of Cu-rich phase precipitation were assumed to allow the Cu precipitation to occur in the auto-aging after acicular ferrite transformation.
The buried pipelines are usually designed to have a lifetime of about 30 to 50 years.1) In order to guarantee this lifetime, pipelines are commonly protected from the external corrosion by a combined application of the highly insulated coating and the cathodic protection to reduce the corrosive attack.2) In spite of these protection measures, cases of failure on underground pipeline due to corrosion, such as microbiologically influenced corrosion (MIC) and hydrogen-induced cracking (HIC) were still continuously reported.3,4,5,6,7,8) Recently, MIC and HIC have been attracted more and more researchers’ attentions because these two problems have caused a great economical loss due to the failure of buried pipelines.1,2) In order to respond to the requirements for pipeline steel to exhibit a good combination of high strength, good toughness balanced with HIC and MIC resistances, a novel low carbon Cu-bearing pipeline steel has been recently developed.9) This development is based on the Cu alloying concept along with a short time of aging following the thermo-mechanical control process (TMCP). Copper (Cu) plays a vital role in this novel steel. Through adding proper amount of Cu to the currently used pipeline steels, fine Cu-rich precipitates with nano-scale could be formed in the steel matrix by adjusting the aging treatment, which could improve strength, ductility, HIC resistance, and anti-MIC performance for steels.9,10) However, improvement of these properties for the novel steel is not an independent factor, which is related to the chemical composition (Cu content), microstructure and technical parameters of thermo-mechanical controlled processing (TMCP) and so on. To balance these properties, it is of significance to develop an optimal microstructure, which can be achieved by precisely designing the chemical compositions and well controlling the TMCP. In this respect, it is necessary to well understand the influence of Cu content on the phase transformation characteristics in the steel during continuous cooling.
It is well known that Cu play an important role in steels, such as exerting a vigorous effect on hardenability,11) enhancing strength via precipitation strengthening,12,13,14) improving fatigue resistance,15) supplying antibacterial activity,10,16) reducing susceptibility of hydrogen embrittlement,17) and lowering ferrite transformation temperature, etc. These indicate that the effect of Cu on microstructure and properties of steels is pronounced. Much research has been done on the effect of Cu on ferrite transformation behavior and related microstructure,18,19,20,21,22,23,24) and it has been already made clear that Cu interphase precipitation could occur only at a certain region of critical cooling rate and temperature.23,24) However, for the novel micro-alloyed Cu-bearing pipeline steel with a small quantity of nickel, chromium, molybdenum and niobium, it is very important to know the effect of Cu addition on the kinetics of austenite decomposition. In this study, therefore, the dynamic continuous cooling transformation (CCT) behavior of three high Cu content pipeline steels was investigated by dilatometry test, optical microscopy observation, transmission electron microscopy observation, and hardness test. The CCT diagrams for different Cu contents were plotted and the microstructural evolution was analyzed. In addition, the precipitation behavior of Cu-rich phase during continuous cooling was also analyzed. These results would lay a foundation for obtaining high performance Cu-bearing pipeline steel with optimized microstructure and mechanical properties.
Three pipeline steels with different Cu contents (1.06 wt.%, 1.46 wt.%, 2.00 wt.%), named 1.0 Cu, 1.5 Cu and 2.0 Cu steel, respectively, were newly designed for experiment. The chemical compositions of the experimental steels are listed in Table 1. They were firstly melted in a 25 kg vacuum induction melting furnace, and the ingots were hot forged to the rods of 30 mm in diameter. Specimens for the compression test were machined from the as-forged rods, with gauge length of 12 mm and diameter of 5 mm, as shown in Fig. 1(a). The specimens were homogenized at 1050°C for 10 min, and then cooled down to the deformation temperature of 980 and 850°C, respectively, at a cooling rate of 5°C/s. After stabilizing holding for 2 s at the deformation temperature, the specimens were deformed to approximately 33% and 38% at the strain rate of 10/s on a Gleeble-3800 thermo-mechanical simulator, and then cooled to room temperature at cooling rates of 0.05, 0.5, 1, 2, 5, 10, 20 and 30°C/s. Figure 2 shows the schematic diagram of thermo-mechanical simulation process.
| Steels | C | Si | Mn | Mo | Cu | Ni | Cr | Nb | S | P | Fe |
|---|---|---|---|---|---|---|---|---|---|---|---|
| 1.0Cu | 0.031 | 0.14 | 1.09 | 0.31 | 1.06 | 0.32 | 0.32 | 0.05 | 0.0011 | 0.005 | Bal. |
| 1.5Cu | 0.019 | 0.12 | 1.03 | 0.31 | 1.46 | 0.31 | 0.31 | 0.05 | 0.0011 | 0.005 | Bal. |
| 2.0Cu | 0.023 | 0.13 | 1.06 | 0.30 | 2.00 | 0.30 | 0.30 | 0.05 | 0.0010 | 0.005 | Bal. |
Schematic diagram of the specimens used in the continuous cooling transformation experiment (a) and section for metallographic samples after compression test (b). The diagram (Fig. 1(b)) on the right hand side shows the sectional view of the 1/4 part of the specimens with the shadow area being the observation position. (unit: mm).
Schematic diagram of thermo-mechanical simulation process.
A combination of optical microscopy, hardness analysis and transmission electron microscopy (TEM) was used to determine the microstructural evolution and precipitation of Cu-rich phase in the experimental steels. Samples for metallographic examination were cut at approximately 1/4 diameter position from the compressed specimens. It was shown in Fig. 1(b). They were mechanically ground to #2000 with sand papers, polished, and then etched in a 2% nital solution. For TEM observation, 300 μm thick disks were cut from the metallographic samples. The disks were firstly mechanically thinned to foils about 50 μm thick and then electro-polished by a twin-jet electropolisher in a solution of 8 vol.% perchloric acid and 92 vol.% ethanol. These foils with very tiny holes were examined by a TEM (FEI Tecnai G2 F20) at an accelerating voltage of 200 kV. Hardness measurements were conducted on polished surface of the metallographic samples using a Vickers hardness tester with a load of 500 g for 15 s, and the average hardness from five different measurements was reported.
Figure 3 shows the typical optical micrographs of the Cu-bearing pipeline steels under different cooling rates. It can be seen that the effects of Cu content and cooling rate on the microstructural evolution were evident. Since the complexity of the microstructure in the ultra-low carbon pipeline steel during the continuous cooling process, its structure constituents are hard to classify.25) In this study, the ferrite microstructures can be mainly divided into polygonal ferrite (PF), quasi-polygonal ferrite (QF), and acicular ferrite (AF), in the orders of decreasing transformation temperatures and increasing cooling rates.26) The polygonal ferrite (PF) and pearlite (P) in microstructure were observed at the cooling rate of 0.05°C/s for three steels, as shown in Figs. 3(a), 3(e) and 3(i). However, 1.0Cu steel shows more pearlite compared with 1.5 Cu and 2.0 Cu steels. It is well-known that carbon is a main element in pearlite. In the present study, unfortunately, the carbon content was not controlled very well during the melting process (Table 1). Higher carbon content increased the fraction of pearlite in 1.0 Cu steel. At the cooling rate of 1°C/s, the transformed microstructure of 1.0 Cu steel was mainly polygonal ferrite, and the pearlite disappeared with increase of cooling rate. However, the grain size of polygonal ferrite was not homogeneous, which is clear to distinguish the grains according to their sizes (Fig. 3(b)). Figure 4 shows the grain size distributions measured by a linear interception method. It can be seen that the grain sizes vary in range of 1–20 μm, 2–17 μm and 2–14 μm for 1.0Cu, 1.5 Cu and 2.0 Cu steeels, respectively. With increase of the Cu content, the polygonal ferrite become homogeneous, and a small number of acicular ferrite was obtained at this cooling rate for 1.5 Cu and 2.0 Cu steels (Figs. 3(f), 3(j)). For further increasing the cooling rate to 5°C/s, a mixture of polygonal ferrite (PF), quasi-polygonal ferrite (QF) and acicular ferrite (AF) were obtained for 1.0Cu steel (Fig. 3(c)). In comparison with that obtained for 1.5 Cu steel, the acicular ferrite dominated at this cooling rate was surrounded by quasi-polygonal ferrite (Fig. 3(g)). The transformed microstructure was full acicular ferrite for 2.0 Cu steel at the cooling rate of 5°C/s (Fig. 3(k)). This microstructure presented various nonequiaxed grain sizes distributed in a random manner.27) Obviously, the microstructures of three steels were transformed to the typical acicular ferrite with increase of the cooling rate up to 30°C/s, as shown in Figs. 3(d), 3(h) and 3(l). Moreover, the average grain size was decreased with increases of the cooling rate and the Cu content. At higher cooling rate, more Cu would be existed as solution state in the steel. The solute Cu atoms in austenite were reported to retard the recrystallization of hot deformed austenite.13) The 2.0 Cu steel retained nonrecrystallized austenite before transformation, which made the acicular ferrite grains more uniform and fine (Fig. 3(l)).
Optical microstructures of experimental steels with different Cu contents after continuous cooling to room temperature under different cooling rates.
The grain size distributions for different Cu-bearing pipeline steels at the cooling rate of 1°C/s: (a) 1.0Cu, (b) 1.5Cu, (c) 2.0Cu.
The continuous cooling transformation (CCT) diagrams are plotted in Fig. 5 under the combination of temperature-dilatation curves and microstructures. As can be seen, all the CCT diagrams are characterized by multilayer transformation curves. The transformation start and finish temperatures rise gently with decreasing the cooling rate. The major transformation curves lie in the temperature range between 780 to 660°C. The features of microstructural evolution in this temperature range include pearlite (P), polygonal ferrite (PF), quasi-polygonal ferrite (QF) and acicular ferrite (AF) when the cooling rate is increased from 0.05 to 30°C/s. For 1.0 Cu steel, the temperatures of Ar3 and Ar1 at cooling rate of 0.05°C/s (equilibrium transformation) are 775 and 666°C, respectively. The full acicular ferrite was produced at cooling rate above 10°C/s (Fig. 5(a)). Comparing the CCT diagram of 1.5 Cu steel with that of 1.0 Cu steel, only a slight difference could be observed. This result seems that the more Cu content has not much influence on the transformation behavior. In general, higher Cu content in the steel is expected to decrease the transformation temperature for austenite to ferrite and then to retard the transformation behavior.22) However, this decrease between 1.0 Cu and 1.5 Cu steels was not obvious in the present study (Fig. 5(b)). As mentioned above, 1.0 Cu steel had more carbon content while less copper content than those in 1.5 Cu steel (Table 1). Carbon and copper are all strong austenite stabilizing elements. The synergistic effect of two elements led to the transformation behavior not changed much for the two steels. However, the temperature range between transformation start and end of 1.5 Cu steel was shrunk at low cooling rate, meaning that the transformation rate got faster at the higher transformation temperature. As for 2.0 Cu steel, CCT diagram had much difference with those of 1.0 Cu and 1.5 Cu steels. It is evident that the transformation temperature for austenite to ferrite was lowered by more Cu addition, and the lowering austenite to ferrite transformation temperature led to an increase in the driving force for the acicular ferrite transformation (Fig. 5(c)). Only at cooling rate above 2°C/s, the full acicular ferrite could be obtained. As same as 1.5 Cu steel, the temperature range between transformation start and end was also shrunk at low cooling rate for 2.0 Cu steel. As mentioned above, the polygonal ferrite and pearlite were produced at the low cooling rate for three steels. Ferrite and pearlite are the diffusional transformation products. The phase transformation is mainly determined by diffusion of carbon atoms. 1.0 Cu steel has higher carbon content, which takes more time to diffuse between austenite and ferrite during phase transition. Accordingly, the temperature range between transformation start and end is expanded. In contrast, 1.5 Cu and 2.0 Cu steels have less carbon content compared with 1.0 Cu steel. Since copper is not a carbide forming element, more copper content has no influence on the interaction between Cu and C atoms. So, the transformation rate is faster, and the temperature range between transformation start and end was shrunk for both 1.5 Cu and 2.0 Cu steels. However, the temperature range between transformation start and end of 1.5 Cu and 2.0 Cu steels was expanded at higher cooling rate, and the transformation start and finish temperatures were shifted to lower temperature. It has been known that acicular ferrite is a transformation product of a mixed diffusion and shear mode.28) At higher cooling rate, the Cu-rich phase has not enough time to precipitate, and more Cu would be presented in the form of solution state, which made the nucleation rate and growth rate of ferrite decreased by the solute drag effect of Cu.22) As a result, the transformation rate of acicular ferrite was retarded and thus the temperature range was expanded.
CCT diagrams of the Cu-bearing pipeline steels: (a) 1.0Cu, (b) 1.5Cu, (c) 2.0Cu.
The effects of Cu content and cooling rate on Vickers hardness of experimental steels are shown in Fig. 6, where there is a strong dependence of HV values on Cu content and cooling rate. As can be seen in Figs. 6(a) and 6(b), the hardness curves of the 1.0 Cu and 1.5 Cu steels behave in a similar manner. The hardness increases with increase of the cooling rate, reaching the highest values of ~215 and ~212 HV for 1.0 Cu and 1.5 Cu steels, respectively. With further increase of the cooling rate, the hardness is not changed much. 1.5 Cu steel has more Cu content compared with 1.0 Cu steel. It is expected that 1.5Cu steel should have higher hardness, but the difference between the two steels is not so large. This is considered to be resulted from the fact that the solid solution strengthening of carbon, just as mentioned above. However, there is a large difference in hardness curve for 2.0 Cu steel. It is evident that a peak is appeared on the curve of 2.0 Cu steels. The hardness curve reaches the highest value of ~254 HV at the cooling rate of 2°C/s. Li et al.23) reported that the precipitation behavior and hardening effect could be affected by cooling rate and copper content in a Fe–Cu–Ni alloy. In the present study, the peak was appeared at cooling rate of 2°C/s for 2.0 Cu steel, not appeared for 1.0 Cu and 1.5 Cu steels. With further increase of the cooling rate, the hardness decreased, and then increased again at cooling rate of 30°C/s. Abe et al.29) reported that the transformation kinetics for austenite to ferrite showed that Cu precipitation occurred only at a certain region of critical cooling rate and temperature in a Cu-bearing high-strength low-alloy (HSLA) steel, which is usually described in terms of interphase precipitation.24) In order to verify this possibility in the present study, the TEM images of 2.0 Cu steel after cooling at 2°C/s were observed. Just as shown in Fig. 7(b), high density particles could be found in the steel matrix. There was no selected area electron diffraction (SAED) pattern to confirm the precipitates because they were so tiny. Luckily, it can be deduced that those spherical particles should be the Cu-rich precipitates according to the chemical compositions of steels (Table 1). Thompson and Liu et al.18,19,20,24) believed that the Cu precipitation took place only in association with polygonal ferrite transformation by interphase precipitation mechanism and no Cu precipitates were detected in the acicular ferrite during continuous cooling. This is consistent with the result of present work, as shown in Figs. 7(a) and 7(c). A larger number of precipitates were dispersed in the polygonal ferrite matrix at a slow cooling rate of 0.05°C/s (Fig. 7(a)). There was no obvious particle observed in the acicular ferrite matrix at a fast cooling rate of 20°C/s (Fig. 7(c)), meaning that the nucleation and growth of Cu-rich precipitates were constrained at higher cooling rate. For the present 2.0 Cu steel at cooling rate of 2°C/s, however, it seems that the result is contradictory. Figure 8 shows the optical and SEM micrographs of 2.0 Cu steel after cooling at 2°C/s, which is clear to see the acicular ferrite microstructure. As mentioned above, the formation of acicular ferrite is a transformation product of a diffusion and shear mixed mode, which allows little time for the diffusion of Cu atoms. So the Cu interphase precipitation should not occur at this cooling rate. However, a large number of nano-scale Cu-rich precipitates were observed in the present study (Fig. 7(b)). Thus, the auto-aging precipitation may be an alternative explanation for the appearance of Cu-rich precipitates. The Cu-rich precipitates occurred during the auto-aging because of the high supersaturation of Cu in austenite (only occur in 2.0Cu steel) and proper cooling rate (enough time for propitiate) in association with acicular ferrite transformation. Seko et al.30) reported that the critical energy required for nucleation was very low in a binary Fe–Cu alloy. Viswanathan et al.31) reported that there was no any incubation period for the onset of Cu-rich precipitation in a 17-4 precipitation hardening stainless steel. Same result was also reported by Mirzadeh et al.32) These simulation predictions and experimental results made it possible for auto-aging precipitation in 2.0 Cu steel at such cooling rate. It indicates that the 2.0 Cu steel had a very short incubation period for the Cu-rich phase precipitation. That is to say that there is no significant free energy barrier for the nucleation process for 2.0 Cu steel. Thus, higher supersaturation of Cu in austenite in 2.0 Cu steel and a short incubation period of Cu-rich phase precipitation allowed the Cu precipitation to occur in the acicular ferrite.
Hardness curves of experimental pipeline steels after cooling at different cooling rates, (a) 1.0Cu, (b) 1.5Cu, (c) 2.0Cu.
TEM images of 2.0Cu steel after cooling at (a) 0.05°C/s, (b) 2°C/s and (c) 20°C/s.
Acicular ferrite microstructures of 2.0Cu steel after cooling at 2°C/s, (a) optical mirograph, (b) SEM mirograph.
(1) The microstructures of novel Cu-bearing pipeline steels formed during continuous cooling consisted of pearlite, polygonal ferrite, quasi-polygonal ferrite and acicular ferrite, depending on the Cu content and cooling rate. A full acicular ferrite could be obtained at cooling rates above 10°C/s, 10°C/s and 2°C/s for 1.0 Cu, 1.5 Cu and 2.0 Cu steels, respectively. More Cu addition could lower the transformation temperature for austenite to ferrite and lead to an increase in the driving force for the acicular ferrite transformation.
(2) The Cu-rich phase precipitation could occur in the acicular ferrite, which made a hardness peak on the hardness vs cooling rate curve of 2.0Cu steel at the cooling rate of 2°C/s. No Cu precipitate was detected in the acicular ferrite at higher cooling rate for 2.0 Cu steel. It is believed that higher supersaturation of Cu in austenite and a short incubation period of Cu-rich phase precipitation allow the Cu precipitation to occur from the auto-aging after acicular ferrite transformation.
The authors would like to acknowledge the financial support by the National Key Technologies R&D Program of China (No. 2011BAE25B03) and National Natural Science Foundation of China (No. 51271175).
