Challenge of mechanical properties of an acicular ferrite pipeline steel

Materials Science and Engineering A 431 (2006) 41–52

Challenge of mechanical properties of an acicular ferrite pipeline steel Fu-Ren Xiao a,b , Bo Liao a,∗ , Yi-Yin Shan b , Gui-Ying Qiao a , Yong Zhong b , Chunling Zhang a , Ke Yang b a

Key Laboratory of Metastable Materials Science and Technology, College of Materials Science and Engineering, Yanshan University, Qinhuangdao 066004, PR China b Institute of Metal Research, Chinese Academy of Science, Shenyang 110016, PR China Received 1 May 2006; received in revised form 5 May 2006; accepted 5 May 2006

Abstract In modern industry, the developing tendency and prospect for productions of the oil and gas pipeline steels is to further improve the strength and toughness by advanced manufacture process of the thermo-mechanical control process (TMCP) to refine microstructure. In this work, the hot deformation behavior as well as its effect on the phase transformation of the clean acicular ferrite pipeline steel with simple chemical composition has been investigated. According to the result, the optimum TMCP parameters were designed. Furthermore, the rolling test was carried out on the experimental rolling mill. The results show that, the high strength and excellent toughness of the clean acicular ferrite pipeline steel can be obtained by controlling the TMCP parameters of the production process appropriately. © 2006 Elsevier B.V. All rights reserved. Keywords: Pipeline steels; Acicular ferrite; TMCP; Process parameters; Mechanical properties

1. Introduction In modern industry, the developing tendency and prospect for productions of the low-carbon micro-alloyed steels is to refine microstructure, so as to further improve its strength and toughness. In order to realize this aim, the thermo-mechanical control process (TMCP) has been applied as an important method. In recent years, great attention has been taken to attain ultra-fine ferrite grain structure by using strain-induced ferritic transformation and/or dynamic recrystallization of ferrite [1,2] during deformation. The microstructure with the finer ferrite grain has been achieved by strain-induced transformation in plain carbon and micro-alloyed steels [3,4]. However, for pipeline steel, if the acicular ferrite microstructure can be achieved, it will with better properties, such as high tensile strength, good toughness, excellent corrosion resistance and superior weldability [5–8], than ultrofine grain ferrite structure [8]. The combination of properties has led to the application of this steel in the manufacturing of large dimension pipes for gas and oil transportation in the low temperature area [9–11].

∗

Corresponding author. Tel.: +86 335 8057047; fax: +86 335 8074545. E-mail addresses: [email protected] (F.-R. Xiao), [email protected] (B. Liao). 0921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2006.05.029

Although the term of acicular ferrite steel was firstly described by Smith et al. in the early 1970s [5,6] and has been widely accepted in pipeline engineering [9–11], there are still controversies and uncertainties on the metallographical identification and classification of the phases. Sometimes, the structure of acicular ferrite is also contested as bainite [12] and quasi-polygonal ferrite (or massive ferrite) [13]. However, from our previous work, the structure of acicular ferrite is considered as the mixture of massive ferrite with bainite ferrite [14]. These results show that the transformation mechanism of the acicular ferrite is different from that of the polygonal ferrite. Meanwhile, when the polygonal ferrite is appeared in the acicular ferrite microstructure, the final mechanical properties of this steel are decreased [8,15], which indicates that the way to refine acicular ferrite structure is different from that with heavy deformation at lower temperature. Recently, the studies on acicular ferrite pipeline steels are focus on the effect of element on the phase transformation and the mechanical properties of acicular ferrite [16,17], however, little systematic work has been carried out in the refinement its microstructure. The acicular ferrite microstructure can be obtained without special alloy elements, and the high tensile strength and good toughness can be obtained too [8,18]. However, the change of the microstructure during hot deformation, and the effect of the hot deformation on phase transformation, as well as the

42

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

refinement of the acicular ferrite of the final production have not been studied extensively. Therefore, in order to achieve the refined acicular ferrite microstructure and to make a technological foundation for this pipeline steel, it is necessary to clarify the changing principle of the microstructure during the TMCP procedures and the effect of the TMCP parameters on the microstructure and mechanical property [19]. In this work, the effect of the hot deformation on static/dynamic recrystallization, strain induced precipitation of carbinitride and phase transformation of the clean pipeline steel (commercial grade X60) with simple chemical composition has been studied. And according to the results, the optimum TMCP parameters were designed. Furthermore, the experimental hot rolling plate of the clean pipeline steel with refined acicular ferrite microstructure was manufactured and its mechanical property was measured by controlling all the design parameters of the production process. 2. Experimental procedure The clean acicular pipeline steel used in this work was prepared in a 25 kg vacuum induction melted furnace, and the cast ingots were forged to bars of Ø25 mm and slabs of 70 mm × 78 mm × 80 mm. The chemical composition of this steel designed by a commercial grade X60 pipeline steel is listed in Table 1. Two types of compressive tests (single-pass and interrupted compressive test) were carried out by means of Gleeble-3500 hot simulator. The hot compressive specimen was machined into Ø10 mm × 15 mm cylinder. For a single-pass compressive test, the specimens were reheated to 1200 ◦ C, then cooled to the testing temperature. The continuous compressive tests were carried out to determine the critical strain in experimental temperature range from 900 to 1100 ◦ C and the strain rate was in range from 0.01 to 10 s−1 . For an interrupted compressive test, the specimens were reheated to 1200 ◦ C and/or 1130 ◦ C, and then cooled to the testing temperature. The interrupted compressive tests were carried out. The experimental technologies were that, the temperature rang was from 900 to 1000 ◦ C, the stain rate was 10 s−1 , the inter-pass time rang was from 1 to 1000 s, and the pre-strain was about 40%. By using the “back extrapolation” method [20,21], the recrystallized fraction (Xa ) after an interval of hot working was calculated. With this method, the softened Table 1 Chemical composition of the experimental steel (wt.%) Element

Content

C Si Mn P S Mo Nb V N O

0.025 0.24 1.56 0.0020 0.0006 0.32 0.039 0.019 0.0062 0.0043

fraction is approximately equal to the recrystallized fraction, in other words, the effect of recovery is excluded from the double compression data [20,21]. The continuous cooling transformation (CCT) curve of this steel with and without hot deformation has been obtained from the experiments by using Formastor-F dilatometer and Gleeble3500 hot simulator. The dimension of the specimen used for the dilatation test was Ø3 mm × 10 mm. The specimens were first heated at 1100 ◦ C for 5 min, subsequently held at 950 ◦ C for 30 s, then cooled to room temperature at linear cooling rates from 0.1 to 100 ◦ C s−1 . The dimension of the specimen used for hot simulative test was Ø10 mm × 15 mm. The specimens were first heated at 1130 ◦ C for 5 min, and then two types of hot deformation processes, the single and two-pass compressive tests, were performed. For the single-pass compression, the specimens were first cooled to 850 ◦ C at a cooling rate of 10 ◦ C s−1 immediately and 40% compressive deformation was carried out at 850 ◦ C, and then the deformed specimens were cooled directly at linear cooling rates from 0.5 to 40 ◦ C s−1 . For two-passes compression, firstly, the specimens were cooled to 980 ◦ C at cooling rate of 10 ◦ C s−1 immediately and 40% compressive deformation was carried out, and then cooled 850 and/or 750 ◦ C at cooling rate of 10 ◦ C s−1 . Secondly, 40% compressive deformation was carried out at same temperature, and then the deformed specimens were cooled directly at linear cooling rates from 0.5 to 40 ◦ C s−1 . A combination of the optical microscopy, dilatometric analysis and transmission electron microscopy (TEM) were used to determine the microstructures of the specimens. The specimens for metallographic examination were mechanically polished and etched with a 3% Nital solution, and then observed by using optical microscopy. For TEM observation, the thin foils were mechanically thinned from 300 to 50 ␮m, and then electropolished by a twin-jet electropolisher in a solution of 10% perchloric acid and 90% acetic acid. The thin foil specimens were observed by using H-800 TEM with 200 kV.

3. Hot deformation behavior of this steel 3.1. Stress–strain curves The typical stress–strain curves of this steel are shown in Fig. 1, which were determined at the temperature range from 900 to 1100 ◦ C and the strain rate range from 0.01 to 10 s−1 . Under the lower temperature or higher strain rate conditions, the stress is increased with the increasing of the strain, whose curve is a typical one of the work hardening. With the increasing of the deformed temperature, or the decreasing of the strain rate, a maximum characterized by peak strain (εp ) and stress (σ p ) is appeared in the curve, followed by a gradual fall to a constant stress value (σ ss ). Furthermore, the peak is changed to narrows and its stress and strain value is decreased. This is typical materials that the recrystallization can be initiated dynamically when the specimen is deformed at temperature above half of their melting point. This is confirmed by an analysis of the effect of the deformation conditions on the peak stress (σ p ) by using the hyperbolic

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

43

Fig. 1. The effect of temperature (a) and strain rates (b) on the stress–strain curves.

sine function [22]: 

ε˙ = A(sinh ασ)n exp




−Qdef RT

3.2. Static recrystallization and strain induced carbonitride precipitation

 (1)

where A, α and n are constants, R is the gas constant, ε˙ the strain rate, Qdef the activation energy, σ the stress and T is the absolute temperature. A method similar to that used by Unira and Jonas [23] was adopted to determine the parameters of the α, n and Qdef at peak stress of dynamic recrystallization. In this case, the activation energy (Qdef ) for deformation was found to be 407 kJ/mol for the tested steel. The results of the fit are presented as a plot of ln(sinh ασ) versus ln Z (Z is the Zener–Hollomon parameter, Z = ε˙ exp(−Qdef /RT )) as shown in Fig. 2. The values of n = 5.9 is in agreement with the observation of other workers [22,24]. Because the strain rates are higher than 10 s−1 or the deformed temperature is below 1000 ◦ C, the Qdef is very high, so that the dynamic recrystallization occurs difficultly. These results indicate that it is difficult to refine austenite grain by utilized dynamic recrystallization during practical industrial rolling process because the strain rate is higher 10 s−1 on rolling mill. Therefore, the static recrystallization will be an effective method to refine initial austenite grain.

The typical interrupted compressive stress–strain curve of the specimens after reheating at 1130 ◦ C are determined at of 950 ◦ C, which are shown in Fig. 3. The experimental conditions are that the constant true strain rate is 10 s−1 , the pre-strain is 40% and the inter-pass times are 1–1000 s. Several interrupted compressive tests with inter-pass time are plotted together to demonstrate the effect of unloading time on the flowing stress–strain curves. It can be seen that, when the inter-pass time is short, the flow curve resulting from second twist follows the continuous one closely because little softening is taken place. With the increasing of the inter-pass time, the initial flow stress of second loading is decreased because static recovery occurs gently. When the inter-pass time is increased to 1000 s, the work hardens of the second curve are similar to that of the first curve so as to rebuild the dislocation structure, which has been recovered because of the static and/or meta-dynamic recrystallization during unloading process [22]. However, it can be found from Fig. 3, when unloading time is in the range from 10 to 120 s, little change can be found from the second curve. It indicates that another factor affects the recovery

Fig. 2. Relationship of ln(sinh(ασ p )) vs. ln Z.

Fig. 3. Double twist stress–strain curves of this steel.

44

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

Fig. 4. Recrystallized fraction plotted against time: (a) austenized at 1130 ◦ C; (b) austenized at 1200 ◦ C.

process. Similar results were obtained on other specimens in different deformation temperature and reheating temperature conditions. The effect of reheating and deformed temperatures on static recrystallization fraction of the steel is shown in Fig. 4, where a plateau can be seen from certain temperatures. This plateau is a consequence of strain induced carbonitride precipitation. The TEM of this precipitation from the selected specimens is shown in Fig. 5. The fine globular precipitates can be observed in the specimens held for 30 s after deformation (Fig. 5a). When the holding time is increased to 1000 s, the number of the fine precipitated globular carbide is increased (Fig. 5b). These precipitated carbide are considered as the Nb(C, N), according to their globular shape [25]. From Fig. 4, it is noticed that the time of 50% softened fraction is delayed obviously when the temperature is below 950 ◦ C. From the plateau exhibited from the soft kinetics curves, it is possible to deduce the start time (Ps ) and finish time (Pf ) of the strain-induced precipitation, which can be used to draw the precipitation–time–temperature (PTT) curves [26]. The PTT curves drawn from the plateau in the soft kinetics curves, are characterized by “C-shape” ones, as shown in Fig. 6. It is noted that the precipitation start time at the lower reheating tempera-

ture were delayed by compared with those at the higher reheating temperature. Recrystallization processes involving nucleation and growth ones can be described by Avrami equation in following way [22]: 
n  t XSRX = 1 − exp −0.693 (2) t50% where XSRX is the fraction of the recrystallized volume and t50% is the time corresponding to half of recrystallized volume, which depends on all variable that intervene in hot deformation, and whose most general expression is described in following way [22]: t50% = Aεp ε˙ q Ds exp

Q RT

(3)

where ε is the strain, ε˙ the strain rate, D the grain size, Q the activation energy, R the gas constant, T the absolute temperature (in K) and p, q and s are constants. The starting time for strain-induced precipitation (tps ) under isothermal conditions can be described by using a model of Dutta and Sellars [27]. The starting time for strain-induced precipita-

Fig. 5. TEM morphology of strain induced carbonitride precipitations at different interrupted time: (a) 30 s; (b) 1000 s.

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

Fig. 6. The PTT curves of this steel.

tion is given for Nb-steels by the following equation [26]: tps = A[Nb]−1 ε−1 Z−0.5 exp

170000 B exp RT T 3 (ln Ks )2

(4)

where Ks = ([Nb][C + 12N/14])/102.26 − 6770/T , Z is the Zener– Hollomon parameter, [Nb], [C] and [N] are the weight percentages of Nb, C and N, respectively, ε is the strain applied during hot rolling and T is the rolling temperature (in K). The effect of strain rate on precipitation times is accounted for

45

in the Zener–Hollomon parameter that is a function of strain rate. From Eqs. (2)–(4), during the same deformation process, the effect of the Nb content in solid solution and the austenite grain size on the starting time for strain-induced precipitation and the activation energy (Q) of static recrystallization is obvious. According to the equation given by Irvine et al. [28], the equilibrium solution temperature of Nb(C, N) carbonitrite is estimated to about 1028 ◦ C. Therefore, it is believed that when the reheating temperature is above 1130 ◦ C, it is enough to dissolve all the Nb(C, N) carbonitrites in austenite. However, the static recrystallization is also affected by primary austenite grain size. With the increasing of the reheating temperature, the size of the primary austenite grain is increased, so that the static recrystallization is restrained. Meanwhile, it also indicates that, by compared with the higher reheating temperature, the precipitation start time is delayed at the lower one. As mentioned above, the effects of the reheating temperature on the static recrystallization and the precipitation for strain-induced Nb carbonitrites, as well as the microstructure refinement and the mechanical properties of final hot rolled products can be found. For the tested steel, the non-recrystallization temperature is 950 ◦ C. Therefore, the reheating temperature, non-recrystallization temperature and carbinitride precipitation should be considerable factors for the production of high performance hot rolled products.

Fig. 7. Effect of hot deformation on CCT diagrams: (a) without hot deformation; (b) with one-pass deformation (40% reduction at 850 ◦ C); (c) with two-pass deformation (40% reduction at 980 ◦ C and 40% reduction at 850 ◦ C); (d) with two-pass deformation (40% reduction at 980 ◦ C and 40% reduction at 750 ◦ C).

46

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

Fig. 8. The optical micrographs of the steel without deformation at cooling rate: (a) 100 ◦ C s−1 ; (b) 10 ◦ C s−1 ; (c) 2 ◦ C s−1 ; (d) 0.2 ◦ C s−1 .

4. Effect of hot deformation on phase transformation The effect of the hot deformation on the phase transformation kinetics and microstructures is obvious [29–31]. Therefore, the effect of hot deformation parameters on CCT curve has been investigated in this work. The CCT curves of this steel with and without hot deformation are shown in Fig. 7, and the microstructures of the specimens at different cooling rates are shown in Figs. 8 and 9, which reveal a pronounced evolution of the microstructures in different cooling conditions.

In CCT curves, the phase transformation microstructures of this steel are complex during the continuous cooling processes, which are bainite ferrite (BF), acicular ferrite (AF), polygonal ferrite (PF), and pearlite (P). The acicular ferrite in this work is neither the titanium oxide one nor welding metals with needlelike grain [32,12]. This nomenclature of the acicular ferrite, first proposed by Smith et al. in the early 1970s [6,7], is readily accepted by the researchers of pipeline steels as an independent microstructure [5,9–11]. This acicular ferrite microstructure is defined as a highly substructured non-equiaxed phase, which is

Fig. 9. Effect of hot deformation on microstructures of cooling rate at 10 ◦ C/s: (a) one-pass; (b) two-pass at 850 ◦ C; (c) two-pass at 750 ◦ C.

600 575 20 22.7 76 750 750 900 899 57 950 957 Designed Measured 26

1050 1052

600 570 20 22.5 76 800 799 920 920 57 950 954 Designed Measured 25

1050 1050

500 509 20 16.1 76 750 750 900 905 57 950 946 Designed Measured 24

1050 1054

500 520 20 17.8 76 800 804 920 922 57 950 955 Designed Measured 23

1050 1047

500 540 20 20 76 750 783 900 900 57 970 969 Designed Measured 22

1100 1080

76 800 800 900 896 57 970 976 Designed Measured

Finish rolling temperature (◦ C)

21

1100 1070

Begin rolling temperature (◦ C) Begin rolling temperature (◦ C)

Conditions Specimen no.

Table 2 TMCP parameters designed and measured during hot rolling

Reduction (%)

Rolling during non-recrystallization Rough rolling

Finish rolling temperature (◦ C)

Reduction (%)

20 18

Cooling rate (◦ C/s)

500

Finish cooling temperature (◦ C)

500

Cooling temperature (◦ C)

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

47

formed during the continuous cooling process by a mixed diffusion and shear transformation mode at a temperature range slightly higher than upper bainite [6,7]. And the characteristic of the acicular ferrite microstructure with bainite in the low carbon Mn–Mo–Nb steels is different from that formed from the primary austenite grain boundary. In the CCT curve without hot deformation, the intermediate transformation such as bainite and acicular ferrite can be promoted by Mn, Mo, Nb elements greatly and the polygonal ferrite and pearlite transformation can be restrained (shown in Fig. 7a). For example, when the cooling rate is changed from 10 to 100 ◦ C s−1 , the microstructure is the classical bainite with clear primary austenite grain boundary (shown in Fig. 8a and b). When the cooling rate is changed from 0.2 to 5 ◦ C s−1 , the acicular ferrite microstructure can be obtained, in which the primary austenite grain boundary is eliminated (shown in Fig. 8c and d). The polygonal ferrite transformation can only be achieved at the cooling rate lower than 0.2 ◦ C s−1 . The effect of hot deformation on CCT curve is very large (shown in Fig. 7b–d). Comparing with the CCT curve without hot deformation (shown in Fig. 7a–d), the bainite transformation region is disappeared with the tested cooling rate, and the acicular and polygonal ferrite transformation curves are moved to the left top corner in the dynamic CCT diagram. Meanwhile, the microstructure shape in the deformed conditions is markedly different from that in the un-deformed ones because of hot deformation. In the transformed microstructure (Fig. 9), whose cooling rate is larger than 10 ◦ C s−1 , the island microstructure is obviously deduced, and the grains are distributed un-homogeneously. Furthermore, the grain and the island are refined. Moreover, the bainite microstructure characteristics with sheaf-like ferrite or primary austenite grain boundary cannot be found at any experimental cooling rates. The characteristics of this microstructure are that of the classical acicular ferrite. It indicates that the effect of the different hot deformed processes on CCT curve and the transformed microstructure is different. The polygonal ferrite/pearlite transformation and acicular ferrite can be promoted by two-pass deformation and the transformed microstructure can be refined remarkably. However, the effect of the second deformed temperature on the beginning formed temperature of the acicular ferrite and transformed microstructure in the experimental condition is small. With the decreasing of the second deformation temperature from 850 to 750 ◦ C, the beginning formed temperature of acicular ferrite is decreased appreciably. When the specimens with one-pass deformation is deformed in the non-recrystallized austenite region, the high density substructure and dislocation are formed in austenite, which increases the nucleation site for the acicular ferrite and promotes the acicular ferrite transformation. However, the specimens with two-pass deformation are first deformed in the recrystallized austenite region and then in the non-recrystallized one. The primary austenitic grain size can be refined by recrystallization during cooling process after first deformation, and then higher density substructure and dislocation are formed in austenite due to finer primary austenitic grain. Therefore, both of refine initial austenitic grain size and deformed austenite can increase the

48

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

Table 3 Mechanical properties of hot rolled plates Specimen no.

Yield strength, σ s (MPa)

Tensile strength, σ b (MPa)

σ s /σ b

Elongation, δ5 (%)

Impact toughness (J, −40 ◦ C)

21 22 23 24 25 26

555 571 588 594 616 626

656 630 650 651 679 716

0.85 0.91 0.90 0.91 0.91 0.87

25 23 24 23 20 23

143 133 149 148 132 133

Note: the impact specimen size is 55 mm × 10 mm × 5 mm.

nucleation site of acicular ferrite and promote the acicular ferrite transformation. 5. TMCP schedules design and microstructure and properties of this steel 5.1. Design of hot rolling process As mentioned above, the effect of all the hot rolling parameters on the final transformed microstructure as well as the final mechanical properties is obvious. Therefore, in order to attain best mechanical properties for the final hot rolling production, the TMCP parameters should be chosen to be optimized. The hot rolling process schedules designed by the results mentioned above were shown in Table 2. The hot rolling experiment was carried out on a pilot rolling mill with Ø370 mm twin rolls. The forged slabs of 70 mm × 78 mm × 80 mm were reheated to 1150 ◦ C for 90 min and rolled to 7 mm by seven passes, then, they were cooled to 500–600 ◦ C at various cooling rates by spraying water on them, and annealed at 500–600 ◦ C for 90 min and cooled in furnace to simulate the coiling process in actual production of pipeline steel plates. The main hot rolling parameters measured from hot rolling experiment are listed in Table 2 too. The higher reheating temperature of 1150 ◦ C is chosen to assure all the alloy element dissolve in austenite and to obtain homogeneous austenite microstructure. The different rough rolling conditions are chosen to obtain the refined and uniformed austenite grain. In order to determine the effect of finish rolling and coiling temperature on the mechanical properties of this steel, they are chosen differently. Considering the effect of deformed reduction on the phase transformation of the acicular ferrite according to the results of CCT curve, the cooling rate is chosen at 20 ◦ C s−1 for ensuring to obtain the acicular ferrite microstructure. 5.2. Mechanical properties and microstructures of rolled plates The mechanical properties of the hot rolling plates by means of various TMCP schedules are shown in Table 3. It can be seen that both extreme high strength and toughness have been obtained. However, the effect of the TMCP parameters on both strength and toughness can be found. The tensile and yield strength of the rolled steel plate are increased with the decreasing of the rough and finish rolling temperature. When the coiling temperature is increased from 500 to 600 ◦ C, the

strength is increased, and the impact toughness and elongation are decreased appreciably. The effect of the TMCP parameters on microstructures of hot rolled plates is shown in Fig. 10. It can be seen that the variety of microstructure is depended upon the TMCP parameters severely. With the decreasing of the rough and the finish rolling temperatures, the microstructure of rolled plate is refined (shown in Fig. 10a–d). With the increasing of finish cooling and/or coiling temperatures, the final average grain size is increased (Fig. 10c–f). However, the acicular ferrite microstructure is confused from the optical metallograph, so the typical specimens were selected to observe by TEM in detailed, which is shown in Fig. 11. From Fig. 11, the microstructure is finer and more complicated, in which the different shape ferrites, such as noneequiaxial and lathy ferrite with dense density dislocations, can be observed. Meanwhile, some martensite/austenite islands are distributed in the microstructure, and a few fine carbonitride particles are distributed in ferrite matrix. This microstructure with fine ferrite grain, martensite/austenite islands, high-density dislocations and fine carbonitride particles is the reason for improving the strength and toughness of this steel. 5.3. Discussion The acicular ferrite microstructure consists of non-equiaxial shape and lathy ferrites with dense dislocations and widely dispersed fine second phases islands in the matrix, because the acicular ferrite transformations consist of two independent phase transformation mode [14,33]. The reason for acicular ferrite pipeline steels with higher properties by the advanced thermomechanical control process is that all the strengthening methods, such as grain refining strengthening, second phase strengthening (both precipitation and M/A island hardening), solution strengthening and dislocation strengthening are utilized in this work [15,34–36]. The advanced thermo-mechanical control process is an effective method to obtain high property by controlling the final acicular ferrite microstructure for micro-alloyed pipeline steels. The hot deformation affects the microstructure refinement by effective nucleation and growth of acicular ferrite, so that, the final microstructure and mechanical properties are depended strongly on controlled rolling parameters and cooling conditions of the plate [37]. Therefore, the controlled rolling parameters (such as reheating temperature, reduction, deforming temperature, interpass time) and cooling conditions (such as cooling rate and finish cooling temperature) play a particularly important role, because

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

49

Fig. 10. Microstructures of rolled plate samples: (a) no. 21; (b) no. 22; (c) no. 23; (d) no. 24; (e) no. 25; (f) no. 26.

the nucleation sites of acicular ferrite can be increased and the acicular ferrite transformation can be promoted effectively by the fine austenite grain and substructure and dislocation with high density formed in austenite [14]. The grain refinement is obtained by control of the rolling conditions—time, temperature and deformations during the whole production process. Grain refinement in the steel is enhanced through a combination of controlled rolling and microalloying. The primary grain refinement mechanism in controlled rolling is the recrystallization of austenite during hot deformation. Small alloying elements like Nb and V can result in the formation of carboni-

trides in the microstructure. These fine precipitates are effective in preventing grain growth. By the use of controlled rolling, recrystallization is retarded during the last pass [38]. From the results of mechanical properties of rolled plates, with the increasing of the rough rolling temperature, both strengths (yield and tensile strengths) and impact toughness is decreased (shown in Table 3). The reason is explained as follow. During rough rolling process, the main effect of processing parameters is on austenite grain size by recrystallization and carbonitrite precipitation. As mentioned above (shown in Figs. 4–6), when the rough rolling temperature is decreased

50

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

Fig. 11. TEM micrographs of sample no. 23.

from the range of 1100–980 ◦ C to that of 1050–950 ◦ C, the fine austenite grain is obtained because the growth of recrystallization austenite grain can be restrained strongly by the strain-induced carbonitride (shown in Figs. 4–6), which results that ultra-fine acicular ferrite microstructures are obtained and both their strength and toughness are improved. These results fit close to those of the microstructure observation. The microstructure is refined with the decreasing of rough rolling temperature (shown in Fig. 10a–d). These results indicate that refinement of primary austenite grain size is an effective method to refine acicular ferrite microstructure and improve the property. On one hand, by decreasing the finish rolling temperature, when the specimen was deformed in the non-recrystallized

austenite region, the substructure and dislocation with the high density were formed in austenite, which increased the nucleation sites of the acicular ferrite and promoted the acicular ferrite transformation, so that it is benefit to refine the microstructure [17]. Furthermore, with the decreasing of the finish rolling temperature from 800 to 750 ◦ C (Fig. 10), the beginning transformed temperature of the acicular ferrite is decreased and the nonequiaxial ferrite size transformed during control cooling process after deformation is smaller than that at high finish rolling temperature. Meanwhile, the density dislocation is increased as a result of the decreasing in phase transformation temperature. The microstructure refinement and high dislocation density result the increasing of strength and toughness of this steel (shown in Table 3).

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

Fig. 12. The different sizes of the precipitated carbonitride of sample no. 26.

From Fig. 10c–f, it can be seen that the coarse microstructure is obtained with the increasing of the finish cooling and coiling temperature. However, its strength is increased, while the toughness is deceased (listed in Table 3 nos. 22–26), which indicates that the properties of this steel will be influenced by other factors besides refinement and dislocation. Two typical carbonitrides with different size distributed in ferrite matrix is shown in Fig. 12. One is larger than 40 nm, which maybe formed during hot rolling process and another is smaller than 10 nm, which is distributed in dislocation and maybe forming during coiling isothermal process. For this steel with Nb and V elements, the peak temperature for carbonitrides of Nb and V precipitation is 600 ◦ C [39,40], so the finish cooling and coiling temperatures are increased from 500 to 600 ◦ C, the amount of precipitated carbonitride is increased. Furthermore, the precipitated carbonitrides are finer because of the high-density dislocation and substructure by hot deformation. The finer carbonitride results in the increasing of the strength and the decreasing of the impact toughness [39]. Therefore, an appropriate coiling temperature is important for precipitation strengthening, which can increase the yield strength and decrease toughness appreciably. As mentioned results above, in order to obtain the optimum mechanical properties with the high strength and excellent toughness for the acicular ferrite pipeline steel, the TMCP parameters during the whole production process must be controlled. All factors of strengthening and toughening, such as the grain refining, second phase strengthening (both precipitation and M/A island hardening), solution strengthening and dislocation strengthening, must be taken into account and utilized comprehensively by controlling the TMCP parameters accurately. In this work, the high strength and excellent toughness of the clean acicular ferrite steel are obtained by controlling the TMCP parameters of the whole production process. 6. Conclusions (1) The activation energy of dynamic recrystallization in Mn–Mo–Nb pipeline steel has been determined. The dynamic and static recrystallization can be restrained greatly and the non-recrystallization temperature can be increased by added Mo and Nb elements into the steel. The effect of the reheating temperature on the behavior of static recrys-

51

tallization and the strain-induced carbonitride precipitation is obvious. (2) The acicular ferrite transformation can be promoted by hot deformation, which makes the CCT curve move toward the top left corner. In addition, the island structures in acicular ferrites become finer, and the grains of the final microstructures are distributed un-homogeneously. (3) The excellent properties of the clean pipeline steel with a simple chemical composition of commercial grade X60 pipeline steel have been achieved grade X80 by controlling all thermo-mechanical control process parameters of the whole production process, which include 580 MPa yield strength, 130 J impact toughness (−40 ◦ C, half size of Charpy impact sample). (4) In order to obtain the good combination of the high strength and toughness for acicular ferrite pipeline steel, it is necessary to control thermo-mechanical control process parameters of the whole production process accurately, which includes the reheating of slabs, controlled rolling, controlled cooling, coiling temperature by comprehensive utilization all strengthening and toughening factors, such as the grain refining strengthening, precipitation hardening, M/A islands second phase strengthening, solution strengthening and dislocation strengthening. Acknowledgments This work is financially supported by “863” project (No. 2005AA412020), the National Natural Sciences Foundation (No. 50471106), Natural Science Foundation of Hebei Province (No. 501205), and Natural Science Foundation of Liaoning Province (no. 20022013). References [1] H. Beladi, G.L. Kelly, A. Shokouhi, P.D. Hodgson, Mater. Sci. Eng. A 371 (2004) 343–352. [2] M. Jahazi, B. Egbal, J. Mater. Process. Technol. 103 (2000) 276–279. [3] S.Y. Ok, J.K. Park, Scripta Mater. 52 (2005) 1111–1116. [4] N.S.V.S. Murty, S. Torizuka, K. Nagai, T. Kitai, Y. Kogo, Scripta Mater. 53 (2005) 763–768. [5] Y.M. Kim, S.K. Kim, Y.J. Lim, N.J. Kim, ISIJ Int. 42 (2002) 1571–1577. [6] Y.E. Smith, A.P. Coldren, R.L. Cryderman, Toward Improved Ductility and Toughness, Climax Molybdenum Company (Japan) Ltd., Tokyo, 1972, pp. 119–142. [7] Y. Smith, A. Coldren, R. Cryderman, Met. Sci. Heat Treat. 18 (1976) 59–65 (English translation of Metallovedenie i Termicheskaya Obrabotka Metallov). [8] M.-C. Zhao, Y.-Y. Shan, F.-R. Xiao, K. Yang, Y.-H. Li, Mater. Lett. 57 (2002) 141–145. [9] Y. Wang, K. Yang, Y. Shan, M. Zhao, B. Qian, Proceedings of the International Pipe Dreamer’s Conference, Yokohama, Japan, November 7–8, 2002, pp. 53–84. [10] T. Janzen, W.N. Horner, Proceedings of the International Pipeline Conference, ASME, New York, 1998, p. 83. [11] Workshops on the Summarization of X70 Grade Pipeline Steel Trial Productions, West-/East Natural Gas Transportation Project Management Organization of Petrochina Co. Ltd., Langfang, China, December 18–21, 2000, p. 3. [12] J.M. Gregg, H.K.D.H. Bhadeshia, Metall. Mater. Trans. 25A (1994) 1603–1612.

52

F.-R. Xiao et al. / Materials Science and Engineering A 431 (2006) 41–52

[13] T. Araki, et al., Continuous-Cooled Zw Microstructures of Low-Carbon Steel, vol. 1, ISIJ, Tokyo, 1992, pp. 4–5. [14] F. Xiao, B. Liao, D. Ren, Y. Shan, K. Ynag, Mater. Charact. 54 (2005) 305–314. [15] T. Tanaka, Int. Met. Rev. 264 (1981) 185–212. [16] J. Kong, L. Zhen, B. Guo, P. Li, A. Wang, C. Xie, Mater. Des. 25 (2004) 723–728. [17] M.-C. Zhao, K. Yang, F.-R. Xiao, Y.-Y. Shan, Mater. Sci. Eng. A355 (2003) 126–136. [18] M.-C. Zhao, K. Yang, Y. Shan, Mater. Sci. Eng. A335 (2002) 14–20. [19] A. Laasroui, J.J. Jonas, Metall. Trans. 22A (1991) 151. [20] S.F. Medina, A. Quispe, ISIJ Int. 41 (2001) 774–781. [21] H.L. Andrade, M.G. Akben, J.J. Jonas, Metall. Trans. 14A (1983) 1967–1977. [22] S.-H. Cho, K.-B. Kang, J.J. Jonas, ISIJ Int. 41 (2001) 766–773. [23] J.L. Uvira, J.J. Jonas, Tans. Metall. Soc. AIME 242 (1968) 1619. [24] A. Laasroui, J.J. Jonas, ISIJ Int. 32 (1991) 95–105. [25] E.V. Pereloma, B.R. Crawford, P.D. Hodgson, Mater. Sci. Eng. A299 (2001) 27–37. [26] S.G. Hong, K.B. Kang, C.G. Park, Scripta Mater. 46 (2002) 163–168. [27] B. Dutta, C.M. Sellars, Mater. Sci. Technol. 3 (1987) 197. [28] K.M. Irvine, F.B. Pickering, T. Gladman, J. Iron Steel Inst. 205 (1967) 161.

[29] J.-C. Zhao, M.R. Notis, Mater. Sci. Eng. R15 (1995) 108–135. [30] V. Homolova, J. Janovec, A. Kroupa, Mater. Sci. Eng. A335 (2002) 290–297. [31] S.C. Hong, S.H. Lim, H.S. Hong, K.J. Lee, D.H. Shin, K.S. Lee, Mater. Sci. Eng. A355 (2003) 241–248. [32] S.S. Babu, H.K.D.H. Bhadeshia, Mater. Sci. Technol. 156 (1992) 1–12. [33] M. Niikura, S. Yamamoto, C. Ouchi, I. Kozasu, ISIJ Int. 70 (1984) 1429. [34] J.M. Gray, in: L. Guoquan, et al. (Eds.), Proceedings of the HSLS Steels’2000, Xi’an, China, 30 October–2 November, Metallurgical Industry Press, Beijing, 2000, pp. 71–79. [35] J. Zrn´yk, T. Kvackaj, A. Pongpaybul, P. Sricharoenchai, J. Vilk, V. Vrchovinsky, Mater. Sci. Eng. A 319–321 (2001) 321–325. [36] A. Bakkalo˘glu, Mater. Lett. 56 (2002) 200–209. [37] T. Siwecki, A. Sandberg, W. Roberts, R. Lagneborg, in: A.J. DeArdo, G.A. Ratz, P.J. Wray (Eds.), Conference of the Proceedings of TMSAIME, Warrendale, USA, 1982, pp. 163–169. [38] F.B. Pickering, Proceedings of the International Conference on Microalloying, New York, 1975, p. 3. [39] M. Charleux, W.J. Poole, M. Militzer, A. Deschamps, Metall. Mater. Trans. A 32 (2001) 1635–1647. [40] H.C. Kang, S.H. Lee, D.H. Shin, K.J. Lee, S.J. Kim, K.S. Lee, Mater. Sci. Forum 449–452 (2004) 49–52.