Novel Cu-bearing economical 21Cr duplex stainless steels

Novel Cu-bearing economical 21Cr duplex stainless steels

Materials and Design 46 (2013) 758–765 Contents lists available at SciVerse ScienceDirect Materials and Design journal homepage: www.elsevier.com/lo...

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Materials and Design 46 (2013) 758–765

Contents lists available at SciVerse ScienceDirect

Materials and Design journal homepage: www.elsevier.com/locate/matdes

Novel Cu-bearing economical 21Cr duplex stainless steels Qingxuan Ran a, Jun Li a, Yulai Xu a, Xueshan Xiao a,⇑, Haifeng Yu b, Laizhu Jiang b a b

Laboratory for Microstructures, Institute of Materials, Shanghai University, Shanghai 200072, China Baoshan Iron & Steel Co., Ltd., Shanghai 200431, China

a r t i c l e

i n f o

Article history: Received 9 September 2012 Accepted 6 November 2012 Available online 23 November 2012 Keywords: Duplex stainless steel Microstructure Mechanical property Corrosion resistance

a b s t r a c t A new family of 21Cr–2 Ni–1.0Mo–0.2 N–xCu (x = 0.5, 1.0, 1.5) economical duplex stainless steels have been developed by examining the effect of Cu on the microstructure and properties of solution-treated specimens. The results have shown that these alloys have a balanced ferrite–austenite duplex structure. The ferrite content increases with the solution treatment temperature, but decreases with an increase in Cu. Some precipitates such as sigma phase, e-Cu and Cr2N were found when solution-treated at 780 °C for 30 min. The yield strength, tensile strength and fracture elongation values of experimental alloys solution-treated at 1020 °C for 30 min were about 540 MPa, 1000 MPa, and 35%, respectively. The pitting corrosion potentials of the solution-treated alloys were all above 500 mV in 1 mol/L NaCl solution at room temperature and the pitting corrosions always occur in ferrite phase. The mechanical properties and corrosion resistance of the designed alloys with lower production cost are better than those of AISI 316L austenitic stainless steel. Ó 2012 Elsevier Ltd. All rights reserved.

1. Introduction The duplex stainless steels (DSSs) are characterized by having a biphasic microstructure, composed of ferrite and austenite with approximately a 1:1 ratio. DSSs have good mechanical properties and corrosion resistance, and, in particular, have excellent pitting resistance in a chloride environment and stress corrosion resistance [1–3]. Due to their excellent mechanical properties and corrosion resistance, DSSs are widely used in the petrochemical, marine, pulp and paper, power industries, etc. [4,5]. Stainless steels are actively being developed with adding copper. Gonzalez et al. [6] has reported the effects of a partial substitution of Ni by Cu on the formability of AISI 304 steel in terms of microstructure changes during tensile straining. They found that the substitution of Ni by Cu gave rise to an increase in the strain level required to induce martensitic transformation, resulting in higher maximum uniform elongation and better stretch formability. Copper is also added in stainless steels as a prominent precipitating element, such as PH 15-5 and 17-4 alloys [7,8]. Cu-bearing antibacterial stainless steels are also developed as a new novel class of structure/function integrated materials in recent years [9–11]. Although Cu has been labeled as a well-known alloying element which is used to improve the corrosion resistance in stainless steels, its effect on the resistance to localized corrosion of stainless steels in chloride media has not been clarified sufficiently. In the ⇑ Corresponding author. Tel./fax: +86 21 56331484. E-mail address: [email protected] (X. Xiao). 0261-3069/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2012.11.017

cases of austenitic stainless steels, Yuan et al. [12] has reported that copper was detrimental to the passivity of stainless steel in acidic 3.5% NaCl solution at 80 °C. But on the other hand, alloying Cu always improved the resistance to the localized corrosion of ferritic stainless steels [13]. Ujiro et al. [14] believed that the beneficial effect of copper in ferritic stainless steels was more than that in austenitic stainless steels. In addition, there are still some controversies between copper’s austenizing power and its effect on corrosion resistance in stainless steels, in particular, in duplex stainless steels. So it is very necessary to investigate the effect of alloying Cu on the properties of duplex stainless steels. The purpose of this work was to develop a new series of economical, Cu-bearing 21Cr duplex stainless steels and investigate the effect of the Cu element on the microstructure and properties of these designed duplex stainless steels. 2. Experimental details The raw material was melted in a 50-kW vacuum induction furnace. The chemical compositions of the castings, represented as Cu-1, Cu-2, and Cu-3, are shown in Table 1. In this experiment, the casting ingots with a diameter of 120 mm were hot forged into bars of 40-mm-outside diameter stick, and the as-forged samples were treated with solution at temperature from 720 °C to 1200 °C for 30 min and then water quenched. The microstructures of the specimens were observed by KEYENCE VHX-100 optical microscopy (Keyence Corporation, Osaka, Japan), and the specimens were electrochemically etched for 45 s at 10 V direct current in 15 wt.% KOH solutions, which made the austenite phase bright

Q. Ran et al. / Materials and Design 46 (2013) 758–765 Table 1 Chemical composition of the experimental stainless steel (wt.%). Steel

C

Cr

Ni

Mo

N

Cu

S

Si

Fe

Cu-1 Cu-2 Cu-3

0.027 0.025 0.029

21.1 20.7 20.4

2.0 2.1 2.9

1.0 1.0 1.1

0.211 0.205 0.207

0.48 0.97 1.50

0.006 0.005 0.005

0.37 0.35 0.41

Bal. Bal. Bal.

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and ferrite phase gray. The ferrite content was measured with quantitative metallographic analysis system. The average of ten measurements on each sample was taken as the ferrite content. X-ray diffraction (XRD) analyses were conducted in a D/max2550 diffractometer (Rigaku, Tokyo, Japan) with Cu Ka radiation (k = 1.5406 Å) and scan rate of 8 deg/min. Tensile tests were made at room temperature with specimens having a gage length of 25 mm and diameter of 5 mm according to the National Standard

Fig. 1. Optical micrographs of experimental alloys solution-treated at 780 °C for 30 min: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

Fig. 2. Optical micrographs of experimental alloys solution-treated at 1020 °C for 30 min: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

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Fig. 3. Optical micrographs of experimental alloys solution-treated at 1200 °C for 30 min: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

Fig. 4. X-ray diffraction analysis of experimental alloys solution-treated at different temperature for 30 min: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

of the PR China GB/T228-2002 [15], and carried out at a strain rate of 1  103 s1. Impact toughness testing was made at 20 °C, 0 °C and 40 °C, respectively; with 10 mm  10 mm  55 mm Charpy

V notched full size specimens, using an AHC-3000/2-AT impact machine with maximum capacity of 450 J according to PR China GB/T229-2007 [16]. Specimens were oriented along the forging

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In order to study the pitting corrosion behavior, corrosion behavior of all samples in 1 mol/L sodium chloride solution was investigated using potentiodynamic measurements at room temperature, at scan rate of 0.5 mV/s. The surface area of the working electrodes was reduced to 1.0 cm2 using epoxy resin, in order to avoid the edge effects. The three compartment electrochemical cell was used, with duplex steel samples as working, a saturated calomel electrode (SCE) as reference while platinum electrode served as counter electrode. Before the scan was initiated, the samples were allowed to remain in the pitting solution for 10 min so as to reach their free corrosion potential. The value of potential exceeding 104 A/cm2 after a sudden increase in electric current was called the pitting corrosion potential. 3. Results and discussion 3.1. Effect of Cu on the microstructure of the solution-treated samples Fig. 5. The ferrite volume fraction of experimental alloys with different solution treatment temperature.

direction with the notch perpendicular to the forging surface. Fracture micrographs after impact tests were observed in a HITACHI SU-1510 scanning electron microscope (SEM). The transmission electron microscope (TEM) specimens were sliced from the solution-treated specimens, ground with abrasive papers and then jet eletropolished with a 90 vol.% anhydrous alcohol and 10 vol.% perchloric acid (HClO4) solution under liquid nitrogen cooling, applying 45 V direct current.

Figs. 1–3 shows the typical optical microstructures of the 21 Cr–2 Ni–1.0Mo–0.2N–xCu samples containing Cu of 0.5%, 1.0%, 1.5%, with different solution treatment temperature from 780 °C to 1200 °C for 30 min followed by water quenching. The microstructures of solution-treated experimental DSSs at the same temperature varied significantly with different copper contents. Under such conditions, the matrix of these alloys showed an obvious duplex structure of austenite and ferrite. The gray etched ferrite (a) islands were embedded in a bright etched austenite (c) matrix and no secondary precipitates were found in the ferrite–ferrite

Fig. 6. The typical TEM graphs taken from specimen, solution-treated at 780 °C for 30 min: (a) sigma phase and e-Cu along phase boundaries, (b) Cr2N nitrides and e-Cu along phase boundaries and (c) e-Cu within phase.

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Fig. 7. The typical TEM images taken from a/c interface of specimen solutiontreated at 1020 °C for 30 min.

Fig. 9. Dislocation structure in ferrite and austenite of Cu-2 solution-treated at 1020 °C for 30 min.

Fig. 8. Tensile results of the experimental alloys solution-treated at 1020 °C for 30 min.

Fig. 10. The Charpy impact toughness of experimental alloys at different temperature.

and ferrite–austenite phase boundaries even at the temperature of 780 °C. The above metallographic results were further confirmed by XRD, which is shown in Fig. 4. The relation of copper to the ferrite volume fraction of experimental alloys with different solution treatment temperature is shown in Fig. 5. It reveals that ferrite volume fraction increased with heat treatment temperature. However, it decreased with the increase of Cu content on the same treatment temperature. By increasing the temperature from 720 °C to 1200 °C, the ferrite volume fractions of Cu-1 and Cu-2 varied from 45% to 60%, while that of Cu-3 varied from 25% to 45%. It seems that Cu acts as an austenite former in this work, and the ferrite content is influenced by both Cu content and solution treatment temperature. The sample containing 1.0 wt.% Cu has a microstructure consisting of ferrite and austenite with approximately a 1:1 ratio in a wide temperature range from 720 °C to 1200 °C. Therefore, the optimal ferrite volume fraction of 50% can easily be obtained by choosing 21Cr–2 Ni–1.0Mo–0.2N–xCu (x = 0.5, 1.0, 1.5) DSSs with 1.0 wt.% Cu content.

Shu et al. [17] has reported that Cu addition and precipitation of

e-Cu phases was detrimental to the pitting corrosion resistance of ferritic stainless steels, it could decrease the pitting corrosion resistance in chloride media. But no secondary phase was found by optical microscopy and X-ray analysis in all specimens treated with solution from 780 °C to 1200 °C for 30 min. As we all know, TEM analysis can show something on such microstructural components which are undetectable with X-ray analysis, so the microstructural features in the specimens solution-treated at different temperature have been studied by transmission electron microscopy (TEM). Fig. 6 shows the typical TEM graphs taken from Cu-2 specimen, solution-treated at 780 °C for 30 min. It presents that e-Cu was observed both along phase boundaries and within phases and some other precipitation such as sigma and Cr2N were also observed along phase boundaries. The size of sigma phase is less than 0.4 lm, that is why the sigma phase is not found in the optical micrograph. It is surprise that e-Cu was observed with different shapes in different areas (spherical-shaped along phase

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Fig. 11. Scanning electron micrographs after impact testing at 40 °C and then electrolytically etched in KOH solution: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

boundaries, spherical-shaped and rod-shaped within phase) and it also appeared with sigma and Cr2N nitrides along phase boundaries. The typical TEM image taken from a/c interface of Cu-2, solution-treated at 1020 °C for 30 min is shown in Fig. 7. It can be seen that there was no precipitate when the solution treatment temperature was 1020 °C. That is to say, experimental alloys have been uniformity when the sample was solution-treated at 1020 °C for

30 min, which is the appropriate solution treatment method for the mechanical and corrosion experiment below. 3.2. Effect of Cu on the mechanical properties The results of the tensile tests carried out at room temperature are presented in Fig. 8. It can be seen that, with the increase of Cu content, the ultimate tensile strength (UTS) decreased first and

Fig. 12. Fractographs after impact testing at 40 °C: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

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Fig. 13. Potentiodynamic polarization curves for experimental alloys in 1 mol/L NaCl at room temperature.

then increased, the elongation increased first and then decreases, and the yield stress (YS) decreased all the time. Both the yield stress (YS) and the ultimate tensile strength (UTS) of 21Cr–2 Ni–1.0Mo–0.2N–xCu (x = 0.5, 1.0, 1.5) duplex stainless steels are higher than those of AISI 316L austenitic stainless steel. There exists a clear relationship between the yield stress (YS) and ferrite volume fraction in these alloys. The Cu-1 presented the highest ferrite volume fraction and the highest yield stress (YS), and the yield stress (YS) of Cu-3 alloy reduced significantly compared to Cu-1 and Cu-2 alloys. The reason of this phenomenon is that the ferrite phase is regarded as stronger than austenite, which make the yield stress (YS) of Cu-1 alloy higher than the other two alloys [18]. But the ultimate tensile strength (UTS) of Cu-3 alloy was even higher than that of Cu-2 alloy. From the dislocation structure in ferrite and austenite of Cu-2 solution-treated at 1020 °C for 30 min in Fig. 9, it is found that the dislocations density

in austenite was much bigger than that in ferrite. As mentioned above, Cu-3 alloy presented the highest austenite volume fraction when solution-treated at 1020 °C for 30 min followed by water quenching. So the strain hardening of Cu-3 alloy made the most contribution to the ultimate tensile strength among the experimental alloys, which made its ultimate tensile strength even higher than Cu-2 alloy. Fig. 10 shows the variation of Charpy impact toughness as a function of experimental temperature with different Cu content. It is clear that the impact energy of these alloys at 20 °C and 0 °C had a large value, but rapidly decreased at 40 °C. At the same temperature, the impact energy of the experimental alloys increased with Cu content in general, but the impact energy of Cu-2 alloy was higher than that of Cu-3 alloy at 40 °C. Fig. 11 shows scanning electron micrographs (SEMs) of experimental alloys after impact test at 40 °C and then electrolytically etching in KOH solution. It is obvious that parts of austenite transformed into martensite at 40 °C. The lath microstructure generated because the phase transformation process from austenite to martensite has caused volume change. The martensite content in Cu-3 was much more than that in Cu-2, which may leads the impact energy of Cu-3 lower than that of Cu-2 at 40 °C. Scanning electron micrographs (Fig. 12) shows detailed fracture morphologies from center regions of the fracture surfaces from Charpy impact specimens determined at 40 °C. For Cu-2 alloy, the ductile fracture mechanism was dominant, as shown in Fig. 12b, where a large number of deep dimples could be observed. For Cu-1 and Cu-3 alloys, the fracture appearance (Fig. 12a and c) revealed small and poorly defined facets connected by tear ridges or shallow dimple, a slight river pattern could be seen within the facets. It suggests that ductile–brittle fracture mechanism has become dominant. The ability to form the martensitic microstructure depends on the martensitic start temperature (Ms), which is primarily composition dependent. In this paper, the martensitic start temperature (Ms) may be estimated by the following equation [19]:

Ms ð CÞ ¼ 1305—41:7Cr—61:1Ni—36:1Mo—5Cu—1667N

Fig. 14. Scanning electron micrographs after polarization testing and then electrolytically etched in KOH solution: (a) Cu-1, (b) Cu-2 and (c) Cu-3.

ð1Þ

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The values of Cu-1, Cu-2 and Cu-3 calculated with formula (1) are 69 °C, 52.66 °C and 61.08 °C, respectively, which are all close to 40 °C. This means that their martensitic transformations may appear when the test was operated at 40 °C, which is agree with the experimental results. 3.3. Effect of Cu on the corrosion properties In order to study the pitting corrosion behavior of experimental alloys, potential dynamic curves were conducted at a scanning rate of 0.5 mV/s in 1 mol/L NaCl aqueous solution at room temperature. Fig. 13 shows the representative polarization curves for experimental alloys with different Cu contents. The pitting corrosion potentials of the experimental alloys decreased from 591 mV to 509 mV with the increase in Cu content from 0.5 to 1.5 wt.%, but all are higher than that of AISI 316L austenitic stainless steel, which is about 470 mV in 3.5% NaCl [18]. It is found that there exists a clear relationship between the pitting potential and content of Cu in these alloys. The Cu-1 presents the lowest Cu content but the highest pitting potential, and the pitting potential of Cu-3 alloy reduces significantly compared to Cu-1 and Cu-2 alloys with the increase of Cu content. This phenomenon in the duplex stainless steels is consistent with the results of ferritic and austenitic stainless steels. The deposited Cu on a corroded surface of the steels could suppress the anodic dissolution of the steels which can improve the resistance to general corrosion of stainless steels, but the stability of deposited Cu decreases in chloride media. It has been reported that the passive film of Cu-bearing stainless steels was destroyed by Cl attack more easily than that of Cu-less steels. The probable reason is that the dissolution of Cu complex ions such 2  2 as CuCl+, CuCl4 , CuCl2 and CuCl3 might assist the destruction of the passive films which is harmful to the pitting corrosion resistance [14,20]. Scanning electron micrographs (SEM) of the samples after polarization tests after electrolytically etching in 10 wt.% KOH solution are shown in Fig. 14. It can be seen that all the pitting holes of experimental alloys occurred in ferrite phase, which indicates that pitting corrosion resistance of austenite phase in these experimental alloys is higher than that of ferrite phase. The pitting corrosion resistance is often correlated with the chemical composition of stainless steels. The pitting resistance equivalent number (PREN) is frequently used to quantify the pitting corrosion resistance of stainless steels. It has been suggested that pitting corrosion depends basically on the content of Cr, Mo and N. According to this hypothesis, pitting resistance equivalent number (PREN) is defined by the following equation [21]:

PREN ¼ %Cr þ 3:3%Mo þ 16%N

ð2Þ

It is well-known that in duplex stainless steel, Cr and Mo enrich in ferrite phase, whereas N is concentrated in austenite phase. So, the pitting corrosion resistance is different between ferrite phase and austenite phase. According to formula (2), the PREN of ferrite phase in Cu-1, Cu-2 and Cu-3 are 24.36, 24.49 and 24.61 respectively, while that of austenite phase are 26.05, 25.93, and 25.80. It is obvious that the PREN of austenite phase are higher than that of ferrite phase, which is basically in accordance with the results of the SEM graphs. 4. Conclusion A new series of Cu-bearing duplex stainless steels 21Cr–2 Ni– 1.0Mo–0.2N–xCu (x = 0.5, 1.0, 1.5) with a balanced ferrite–austenite relation, better mechanical and corrosion property have been

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prepared and characterized. With the increase of Cu content, the ferrite volume fractions, yield strength (YS) and pitting corrosion potential decrease, the ultimate tensile strength (UTS) decreases first and then increases. The Charpy impact toughness increase first and then decrease at 40 °C. The reason is that parts of austenite have transformed into martensite. Some precipitates, such as sigma phase, e-Cu and Cr2N, are found when solution-treated at 780 °C for 30 min. 21Cr–2 Ni–1.0Mo–0.2N–1.0Cu is found to have a excellent comprehensive properties with a better combination of UTS, YS, elongation and corrosion resistance compared with those of other alloys. All of these designed alloys with a lower production cost have a higher mechanical strength and corrosion resistance than those of AISI 316L austenitic stainless steel. Acknowledgements The authors would like to thank J.C. Peng for the excellent support during JEM 2010F TEM observations, B. Lu for XRD analysis and Y.L. Chu for SEM analysis in Instrumental Analysis & Research Center of Shanghai University. The research was supported by the National Key Technology R&D Program of China (2012BAE04B02). References [1] Kacar R. Effect of solidification mode and morphology of microstructure on the hydrogen content of duplex stainless steel weld metal. Mater Des 2004;25: 1–9. [2] Nilsson JO. Super duplex stainless steels. Mater Sci Technol 1992;8:685–700. [3] Farnoush H, Momeni A, Dehghani K, et al. Hot deformation characteristics of 2205 duplex stainless steel based on the behavior of constituent phases. Mater Des 2010;31:220–6. [4] Fréchard S, Martin F, Clement C, Cousty J. AFM and EBSD combined studies of plastic deformation in a duplex stainless steel. Mater Sci Eng A 2006;418:312–9. [5] Sato YS, Nelson TW, Sterling CJ, Steel RJ, Pettersson CO. Microstructure and mechanical properties of friction stir welded SAF 2507 super duplex stainless steel. Mater Sci Eng A 2005;397:376–84. [6] Gonzalez BM, Castro CSB, Buono VTL, Vilela JMC, Andrade MS, Moraes JMD, et al. The influence of copper addition on the formability of ALSI 304 stainless steel. Mater Sci Eng A 2003;343:51–6. [7] Habibi Bajguirani HR. The effect of ageing upon the microstructure and mechanical properties of type 15-5 PH stainless steel. Mater Sci Eng A 2002;338:142–59. [8] Hsiao CN, Chiou CS, Yang JR. Aging reactions in a 17-4 PH stainless steel. Mater Chem Phys 2002;74:134–42. [9] Ren L, Li N, Yang K. Study of copper precipitation behavior in a Cu-bearing austenitic antibacterial stainless steel. Mater Des 2011;32:2374–9. [10] Zhang ZX, Lin G, Jiang LZ, Xu Z. Beneficial effects of nitrogen on austenite antibacterial stainless steels. Mater Chem Phys 2008;111:238–43. [11] Li N, Liu Y, Lü M. Study on antibacterial mechanism of copper-bearing austenitic antibacterial stainless steel by atomic force microscopy. J Mater Sci 2008;19:3057–62. [12] Yuan J, Wen L, Shen W. The effect of copper on the anodic dissolution behavior of austenitic stainless steel in acidic chloride solution. Corros Sci 1992;33:851–9. [13] Seo M, Hultquist G, Leygraf C, Sato N. The influence of minor alloying elements (Nb, Ti and Cu) on the corrosion resistivity of ferritic stainless steel in sulfuric acid solution 1986;26:957–60. [14] Ujiro T, Satoh S, Staehle RW, Smyrl WH. Effect of alloying Cu on the corrosion resistance of stainless steels in chloride media. Corro Sci 2001;43:2185–200. [15] GB/T 228-2002. Metallic materials-tensile testing at ambient temperature. China Standard Press; 2002. [16] GB/T 229-2007. Metallic materials-Charpy pendulum impact test method. China Standard Press; 2007. [17] Shu J, Bi HY, Li X, Xu Z. The effect of copper and molybdenum on pitting corrosion and stress corrosion cracking behavior of ultra-pure ferritic stainless steels. Corros Sci 2012;57:89–98. [18] Wang J, Uggowitzer PJ, Magdowski R, Speidel MO. Nickel-free duplex stainless steels. Scripta Mater 1999;40:123–9. [19] Lo KH, Shek CH, Lai JKL. Recent developments in stainless steels. Mater Sci Eng R 2009;65:39–104. [20] Lin HT, Tsai WT, Lee JT, Huang CS. The electrochemical and corrosion behavior of austenitic stainless steel containing Cu. Corros Sci 1992;33:691–7. [21] Baba H, Kodama T, Katada Y. Role of nitrogen on the corrosion behavior of austenitic stainless steels. Corros Sci 2002;44:2393–407.