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Effect of N on microstructure and mechanical properties of 16Cr5Ni1Mo martensitic stainless steel
Materials and Design 34 (2012) 74–81
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Effect of N on microstructure and mechanical properties of 16Cr5Ni1Mo martensitic stainless steel X.P. Ma a, L.J. Wang a, B. Qin b, C.M. Liu a,⇑, S.V. Subramanian c a
Key Laboratory for Anisotropy and Texture of Materials, Northeastern University, Shenyang 110004, China Research Institute for Stainless Steel, R&D Center, Baoshan Iron & Steel Co., Ltd., Shanghai 201900, China c Department of Materials Science and Engineering, McMaster University, Hamilton, Canada L8S-4L7 b
a r t i c l e
i n f o
Article history: Received 26 May 2011 Accepted 27 July 2011 Available online 3 August 2011 Keywords: A. Ferrous metals and alloys E. Mechanical F. Microstructure
a b s t r a c t The effect of adding N (about 0.1% in mass) to low carbon 16Cr5Ni1Mo martensitic stainless steel by adding nitride ferroalloy under argon protective atmospheric of near normal pressure on phase transformation and microstructure of steels subjected to normalizing and tempering was investigated by using dilatometer, laser scanning confocal microscopy, transmission electron microscopy and X-ray diffraction, and its consequence on mechanical properties was evaluated by tensile and impact tests. N can effectively avoid the occurrence of d ferrite during solidification. N retards the kinetics of the formation of reversed austenite and promotes the occurrence of Cr2N during tempering above 550 °C. The strength properties of the steel as normalized is enhanced by N alloying without much decrease in elongation and toughness due to the combined effects of N solution strengthening, suppression of d ferrite and retained austenite. The slight influence of N addition on mechanical properties of martensitic stainless steel after tempering above 550 °C is attributed to the precipitation of coarse rod-like Cr2N and loss of coherence. While the increase in volume fraction of retained austenite originated from reversed austenite with tempering temperature contributes to the decrease in strength properties and increase in toughness and elongation, the retransformation of reversed austenite into martensite affects these properties inversely. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction Super low carbon martensitic stainless steels called super martensitic stainless steel (SMSS) are currently produced for the applications of oil country tubular goods (OCTG) [1]. They are based on the Fe–Cr–Ni–Mo system with 13–16 mass% Cr, 4–6 mass% Ni, 0.5– 2.5 mass% Mo. The basic design concept is the increase in effective Cr content by reducing C to improve corrosion resistance and weldability, the addition of Ni to maintain martensitic microstructure without d ferrite upon cooling and to achieve reversed austenite during tempering, and the addition of Mo to improve the resistance to localized corrosion and sulfide stress cracking (SSC). Martensitic microstructure can be achieved by air cooling from austenite single field due to their adequate hardenability. Retained austenite originated from reversed austenite during tempering is very effective in contributing to the high toughness and ductility of SMSS, which is another advantage of SMSS besides corrosion resistance. There has been a commercial 16Cr5Ni1Mo martensitic stainless steel. Higher chromium content of 16% in the commercial steel is designed to slow down the overall corrosion rate of the steel. How⇑ Corresponding author. Tel.: +86 24 83687554; fax: +86 24 23890920. E-mail address: [email protected] (C.M. Liu). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.07.064
ever, Cr is a ferrite former and an excessive Cr content causes d ferrite to be generated, which not only impairs the toughness properties of the steel but also can generate SCC originated from the inter-phase between the martensite and ferrite [2–4]. Once the d ferrite forms, it is difficult to remove it by conventional heat treatments (except for high temperature diffusion annealing [5,6], which causes burning loss and severe distortion of castings). Thereby, it is necessary to make alloying additions so that austenite phase is stabilized in order for avoiding d ferrite and achievement of more retained austenite along martensitic boundaries and within laths during tempering to get toughness. Ni is the common alloying element used to stabilize the austenite or expand the c phase field, but more Ni application will add on cost. According to Delong and Reid diagram [7], the nickel equivalent number has to be increased by three to balance chromium equivalent number of 17 in commercial 16Cr–5Ni–1Mo martensitic stainless steel to avoid the formation of d ferrite. N is a cost effective austenite stabilizer, which is about 25–30 times more effective than Ni [7], and about 0.1 wt% addition of N is expected to effectively avoid the occurrence of d ferrite in 16Cr–5Ni–1Mo martensitic stainless steel. However, the low solubility of N in martensite make the additions of N in martensitic stainless steel difficult, and several special routes for production of high nitrogen martensitic stainless
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to the thickness of 100 lm and electropolished with a double jet Tenupol Struers machine at 40 °C, 15 V in an electrolyte containing 27 ml perchloric acid and 273 ml alcohol. Carbon extraction replica were prepared on sample surface well-polished and lightly etched using electrolyte containing 2 ml HCl + 4 ml HNO3 + 12 ml alcohol. The precipitate morphology, size and dispersion were examined using conventional transmission electron microscopy (TEM-Philips CM12 120 kV). The micro chemical compositions of precipitates were analyzed using Energy Dispersive Spectrometer (EDS) attached to TEM. The volume fraction of the retained austenite was measured at room temperature by X-ray diffraction using Cu Ka radiation on tempered specimens. The evaluation of volume fraction was determined by measuring the integrated intensities of (1 1 1)c and (1 1 0)a peaks of X-ray pattern, using the procedure outlined by Leem et al. [18].
steels has been used, such as pressurized-electro-slag-remelting [8], powder metallurgy and thermo-chemical treatments [9]. Meanwhile, a lot of works has been done to study the effect of introduction of N on microstructure and mechanical properties of martensitic stainless steel. In general, these works mainly aimed to combine elevated hardness and wear resistance with good corrosion resistance by substitution of C by N through more homogeneous precipitation of finely dispersed Cr nitride in Fe–Cr–N and Fe–Cr–Mo–N alloys [8–13]. Calliari et al. [14] has also studied the microstructure and properties of a martensitic stainless steel alloyed with 0.08 mass% N, but no specific effect of N was concerned. The present study tentatively use N to partially replace Ni. The effect of introduction of N into commercial super low carbon 16Cr– 5Ni–1Mo martensitic stainless steel under argon protective atmosphere on microstructure evolution during heat treatment, and its consequence on mechanical properties including tensile and Chary impact toughness were investigated to generate base line information for high nitrogen Fe–Cr–Ni–Mo martensitic stainless steel.
3. Results and discussion 3.1. Microstructure characterization
2. Experimental procedure Fig. 1 shows the LSCM images of LN and HN2 steels as normalized disclosing martensitic microstructure achieved by air cooling the steels from austenitizing temperature of 1100 °C. Prior austenite grains were subdivided into packets of martensite laths. About 6.5% d ferrite occurs along prior austenite grain boundaries of LN steel due to its higher Cr equivalent. Compared with LN steel, the amount of d ferrite is significantly decreased to 1.2% and 0.7% in HN1 and HN2 steels, respectively by introduction of 0.12 wt% and 0.096 wt% N to LN steel due to strong austenite stabilizing effect of effect N. This is in agreement with original design concept. N can be used to replace Ni to avoid d ferrite effectively. Careful TEM observation on as normalized samples with high nitrogen did not reveal the presence of N rich precipitates. It is supposed that N was entirely dissolved into austenite during solution treatment at 1100 °C and did not precipitate during air cooling. Fig. 2 shows the TEM image of detailed microstructure of LN steel as normalized showing highly dislocated martensitic laths of 0.2–0.3 lm thick and d ferrite. N as very strong austenite stabilizer may increase the volume fraction of retained austenite after normalizing, which will increase the toughness of the steel, but decrease the strength. However, X-ray diffraction test could not prove this point due to low volume fraction of retained austenite in as normalized samples. The comparison of X-ray diffraction pattern of three steel samples as normalized did not found apparent influence of N on lattice parameter, and the structure of N alloyed martensite remains body-centered cubic, other than tetragonal. Fig. 3 shows the austenite transformation volume fraction as a function of heating temperature calculated from the continuous heating curve of dilation diagrams of as-normalized LN and HN2 steel samples at heating rate of 3 and 60 °C/min, using the procedure outlined by Lee et al. [19]. As can be seen, the curves of the same steel at higher heating rate are shifted towards higher temperatures, which means that transformation involved in the austenitization process is heating rate sensitive. At lower heating rate, the phase transformation temperature will get closer to equilibrium point. The curves are shifted towards high temperature for HN2 steel compared with LN steel, demonstrating that addition of
High nitrogen steel samples were made in laboratory in the form of ingots by melting in a 100 kg vacuum induction furnace under argon protective atmosphere of near normal pressure through adding nitride ferroalloy (nitrogen of 5%) and electroslag remelting under atmospheric pressure [15]. The chemical compositions of three tested steels with and without N are summarized in Table 1. The content of N was determined using the spark emission ARL 4460 spectrometer. While commercial LN steel is the control steel without N, HN1 and HN2 are indicated as high nitrogen stainless steel. The ingots were hot rolled at 1200 °C into plates with 12 mm thickness. Normalizing was carried out at 1100 °C for 0.5 h, followed by tempering for 2 h at temperatures from 550 °C to 700 °C. After tempering, the samples were quenched in oil. The heating rate of tempering treatment is about 60 °C/min. The dilatometer technique was used to study the thermal expansion of the steels during continuous heating treatments, which were carried out by heating as-normalized samples from room temperature at different rate of 3 °C/min and 60 °C/min to the austenitizing temperature of 1100 °C. The mechanical properties were evaluated by means of tensile and impact tests. The tensile tests were performed on a tensile testing machine CMT5105 with a extensometer at strain rate of 0.001 s 1. Cylindrical test specimens with a uniform gauge cross-section of U 5 25 mm and machined parallel to the plate rolling direction were used [16]. The impact tests were performed on full size Charpy V notch impact specimens (10 10 55 mm) at room temperature [17]. All mechanical properties value reported in the results are the average of three tests. Metallographic specimens were cut from broken impact samples. The microstructures of samples were investigated by using Laser scanning confocal microscopy (LSCM) and transmission electron microscopy (TEM). The fraction of d ferrite was estimated by quantitative image analysis software at 200 magnification according to the contrast between d ferrite and martensite matrix, and the final value reported is the average of five measurements. Thin foils for TEM analysis were prepared by mechanically grinding the samples with initial thickness of 400 lm
Table 1 Chemical compositions of the tested steels (in mass%). Steel grade
C
Si
Mn
P
S
Cr
Ni
Mo
N
Fe
LN HN1 HN2
0.03 0.017 0.015
0.32 0.32 0.32
0.65 0.78 0.8
0.011 0.011 0.011
0.011 0.007 0.008
16.37 16.32 16.26
5.12 5.03 5.01
0.83 0.98 1.03
0.0079 0.096 0.12
Bal. Bal. Bal.
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Fig. 1. LSCM images of (a) LN and (b) HN2 steels as normalized showing martensitic microstructure and d ferrite distributed along prior austenite grain boundaries.
Fig. 2. TEM image of LN steel as normalized from austenitizing temperature of 1100 °C showing d ferrite and highly dislocated martensitic laths of 0.2–0.3 lm in thickness.
Fig. 3. The variation of austenite volume fraction with heating temperature calculated from continuous heating dilation diagram of two SMSSs with heating rate of 3 and 60 °C/min.
N retards the austenite transformation. This is contrary with the original purpose of design of high nitrogen martensitic stainless steel. N as strong austenite stabilizing element was expected to promote the formation of reversed austenite. According to the
work by Kaluba et al. [20] on martensitic stainless steel with nitrogen, the process of austenitizing of martensite during heat treatment is mainly affected by diffusion of interstitial element C and N. It is proposed in present study that the kinetics of austenite transformation is restrained due to the additional diffusion of N. As determined in Fig. 3, although the Ac1 temperature is further above 550 °C up to 625 °C and 850 °C at heating rate of 3 °C/min and 60 °C/min, respectively, the reverse transformation from martensite to austenite would occur if the samples were maintained above 550 °C in present study. In agreement with many studies carried out on microstructure evolution of martensitic stainless steel during tempering [21,22], reversed austenite obtained by tempering the martensite is relatively stable upon cooling to room temperature and can be partially retained at room temperature. Previously published work by Qin et al. [23] on LN steel has shown that the volume fraction of retained austenite in samples tempered at various temperatures for 2 h increases with increasing tempering temperature and peaks after tempering at 600 °C, above which the volume fraction of retained austenite decreases due to the retransformation of reversed austenite to martensite owing to lowered stability of reversed austenite after tempering at higher temperature. It has been reported that the stability of reversed austenite during cooling after tempering relates to its chemical compositions as well as tempering temperatures [24]. The concentration of alloy elements in reversed austenite during tempering makes it stable due to the decreased MS. With higher tempering temperature, the quenched-in vacancies are expected to increase, which would lower the stability of reversed austenite to be retained upon cooling. Most of reversed austenite can be retained after tempering at lower temperature due to its higher stability. As the temperature is increased further, the reversed austenite volume fraction increases rapidly so that it is only slightly enriched in alloy content and largely retransformed into ‘‘fresh’’ martensite on cooling. Fig. 4a shows the results of X-ray diffraction carried out using Cu Ka radiation on samples of three steels tempered at 600 °C for 6 h. According to the calculation procedure outline by Leem [18], the volume fraction of retained austenite is about 14.4%, 9.3%, 7.2% in LN, HN1 and HN2 steel, respectively. Due to the retardation effect of N on formation of reversed austenite, the tempering temperature for certain volume fraction of reversed austenite, at which the concentration of alloying elements in reversed austenite decreases and composition of reversed austenite is close to average contents of the steel, is higher for HN steels than for the LN steel, thus the retransformation of reversed austenite to martensite would occur at different extent after tempering at same temperature, more in LN steel. As shown in Fig. 4b, while the peaks of retained austenite are present in X-ray diffraction pattern of HN steel tempered at 650 °C for 2 h, the diffraction pattern on LN steel
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Fig. 4. X-ray diffraction results obtained by Cu Ka radiation on samples tempered at (a) 600 °C for 6 h and (b) 650 °C for 2 h.
sample tempered at same temperature for 2 h shows no retained austenite peak, demonstrating that the volume fraction of austenite retained in HN2 steel is higher than that in LN steel after tempering at 650 °C. Fig. 5a and b exhibit TEM bright field and dark field pictures of retained austenite of 100–200 nm in thickness in HNl steel after tempering at temperature of 600 °C. Fig. 5c and d show TEM morphologies of retained austenite in HN1 steel after tempering at 700 °C using bright field and dark field, respectively. The retained austenite film is found with thickness of about 50 nm. Due to the limitation of equipment, the segregation of N between retained austenite and matrix can not be exactly detected.
Previous studies [2,25] have shown that the precipitates in low carbon martensitic stainless steel during tempering are mainly M23C6, where M represents Cr, Mo and Fe. It is reported that the addition of nitrogen delays carbide precipitation and results in disappearance of carbide phase in high nitrogen martensitic stainless steel [26]. Previous work [27] by the authors on microstructures evolution of 15Cr martensitic stainless steel with 0.16 wt% N during tempering has shown that while tempering below 500 °C does not show Cr2N precipitation under TEM observation, tempering above 500 °C results in severe precipitation of Cr2N along martensitic inter-lath boundaries and within matrix. Although the solubility limit of N in martensitic stainless steel is unknown, recent study
Fig. 5. (a) and (b) Bright field and dark field TEM images of HN1 sample tempered at 600 °C showing retained austenite with thickness of 100–200 nm distributed along lath boundaries and within laths; (c) and (d) bright field and dark field TEM images of HN1 sample tempered at 700 °C showing martensite laths and retained austenite with thickness of 50 nm.
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found that the thermodynamic potential for precipitation of Cr2N even occur in 13Cr–5Ni–1Mo martensitic stainless steel with only 0.04 wt% N addition. Fig. 6 shows the TEM images of morphologies, dispersion of precipitates observed in HN1 steel tempered at various temperatures. Fig. 6a shows the TEM pictures of rod-like pre-
cipitates occurring at martensitic lath boundaries and within the lath of HN1 steel tempered at 550 °C. Fig. 6b and c show the TEM pictures of precipitates within martensite lath and at lath boundaries in HN1 steel tempered at 600 °C, respectively. The rod-like precipitates within matrix is around 10 nm in diameter and
Fig. 6. TEM pictures of HN1 steel tempered at: (a) 550 °C showing precipitates along lath boundary and within the lath; (b) and (c) 600 °C showing precipitates distributed randomly within matrix and along lath boundaries; (d) and (e) TEM pictures of carbon extraction replica of HN1 steel tempered at 600 °C showing precipitates well aligned in arrays and randomly distributed, respectively; (f) TEM images of thin foil of HN1 steel tempered at 700 °C showing boundary precipitates of 15 nm in diameter and 30–80 nm in length; (g) ED spectrum analysis of precipitates exhibited in (d) and (e) showing X-ray signals characteristic of Cr and small amount of Mo; (h) SADP of precipitates shown in (e).
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50 nm in length, and precipitates at lath boundaries is continuous. Due to low contrast between precipitates and matrix, and the interference from matrix, it is difficult to observe and identify the precipitates using thin foil. Fig. 6d and e show TEM pictures of carbon extraction replica of HN1 steel tempered at 600 °C showing precipitates well aligned in arrays with inter-spacing comparable with inter-lath spacing of martensitic lath shown in Fig. 6c and precipitates randomly dispersed, respectively. It is considered that the precipitates well aligned were nucleated at boundaries and the precipitates randomly dispersed were nucleated within martensitic laths. The ED spectrum of all precipitates shows X-ray signals characteristic of Cr, Mo and Fe, as shown in Fig. 6g, where Cu arises from copper grid. From the selected area diffraction pattern (SADP) shown in Fig. 6h, these precipitates are identified with hcp structure with parameters of a = 0.480 nm, c = 0.447 nm. According to the crystalline structure and lattice parameter as well as the EDS analysis, these precipitates are identified as Cr2N. It is suspected that significant depletion of Cr in vicinity of Cr rich precipitates at lath boundaries would occur due to high diffusion rate of substantial element at boundaries, which will deteriorate corrosion resistance. Fig. 6f shows TEM picture of HN1 steel tempered at 700 °C showing Cr2N of 15 nm in diameter and 30–80 nm in length. With increasing tempering temperature, Cr2N grew slightly. Actually, the precipitation of Cr2N in this temperature range is undesirable. However, in order to get more retained austenite to restore the ductility and toughness, it is necessary to tempering the steel above 550 °C at which Cr2N would precipitate simultaneously. 3.2. Mechanical properties Fig. 7 shows the tensile properties, Charpy impact toughness of three tested steels at room temperature after normalizing and tem-
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pering at temperatures varied from 550 °C to 700 °C for 2 h. The yield strength is 775, 795, 815 MPa, and the ultimate tensile strength is 1005, 1055 and 1115 MPa for LN, HN1 and HN2 steels, respectively in normalized condition. It is noted that while the yield strength and ultimate tensile strength of the steels as normalized is enhanced to some extent by N addition in two HN steels, the elongation and Charpy impact toughness is slightly improved for HN1 and decreased for HN2 steel. Bahrami and Hendry [9] reported that the strength increases linearly with increasing nitrogen content in 10% Cr martensitic stainless steel in solution treated condition. This is inconsistent with the results in present study. Combined with comparison of microstructure of the steels with and without N presented above, the difference in mechanical properties between steels with and without N after normalization is not only attributed to the solution strengthening effect of N, but also suppression of d ferrite and possible increase in volume fraction of retained austenite by N. Wang et al. [2] found d ferrite is soft owing to its relative low C content and dislocation density in low carbon martensitic stainless steel. Suppression of d ferrite plus solution strengthening effect of N is supposed to significantly increase the strength of the N alloyed martensitic stainless steel. However, the increment in strength caused by N addition is much lower than that reported in literatures. In addition, Carrouge et al. [3] and Wang consistently proposed that the presence of d ferrite in low C martensitic stainless steel does not change the upper and lower shelf energy apparently, while lowers the impact energy remarkably in the transition temperature range and raises the ductile to brittle transition temperature. As the temperature of 15 °C employed in testing the Charpy impact toughness in present study is above the transition temperature range reported in literatures, the suppression of d ferrite is not considered to be the main factor affecting the tensile properties and impact toughness at room tem-
Fig. 7. Mechanical properties of the three studied steels as a function of tempering temperature at room temperature. (a) Yield strength (Rp0.2). (b) Ultimate tensile strength (Rm). (c) Elongation (A%) and (d) Charpy impact toughness (Akv).
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perature. Combining above effects of N on tensile strength properties and Chary toughness, it is deducted that the increase in retained austenite volume fraction in N alloyed martensitic stainless steel in as normalized condition significantly contributes to the improved toughness and slightly increased strength of N alloyed steels in comparison with the control alloy without N. While tempering above 550 °C promotes the precipitation of Cr2N within martensitic laths and at martensitic lath boundaries, strengthening the alloy at cost of elongation and toughness, retained austenite is formed along lath boundaries and within laths simultaneously, restoring ductility and toughness. The mechanical properties of the tested steels after tempering depend on the balance among softening effect of retained austenite, recovery effect of martensite matrix, and the hardening effect of precipitates and the retransformation of reversed austenite to martensite. As shown in Fig. 7a, the yield strength of three tested steels is increased by tempering at 550 °C in relation to as normalized condition due to the precipitation of Cr carbide in LN steel and Cr nitride in HN steels. The ductility and toughness are effectively improved by tempering treatment, owing to the recovery of martensite matrix and formation of small amount of retained austenite during tempering, as shown in Fig. 7c and d. As tempering temperature is increased, the yield strength and ultimate tensile strength first decrease to minimum values after tempering at 625 °C, 650 °C and 600 °C, 625 °C for LN and HN steels, respectively, and then revert to increase, while the elongation and toughness change contrarily, first increasing to maximum values at 600 °C, 625 °C and 600 °C, 650 °C, respectively and then revert to decrease. Retained austenite originated from reversed austenite is reported largely to be very effective in increasing toughness and elongation. The increase in volume fraction of retained austenite apparently improves ductility and toughness, and its decrease lowers both ductility and toughness. This is also can be confirmed in present study. Combined with the measurement of retained austenite volume fraction in LN steel presented previously, the temperature at which the elongation and toughness of LN steel are highest is in agreement with the temperature at which the highest volume fraction of retained austenite has been registered. Deducing from this point, the increase in strength properties and the decrease in ductility and toughness of HN steels after tempering at higher temperature are mainly attributed to the retransformation of reversed austenite to martensite. On the other hand, the retardation of N on formation of reversed austenite would retard the softening of martensite during tempering. Comparing all mechanical properties of three tested steels tempered in the temperature range from 550 °C to 700 °C, the strength properties is slightly enhanced, while the Charpy toughness of steel is severely decreased by N alloying due to the precipitation of Cr2N at lath boundaries and within matrix. Previous work by other researchers [9,12] and the author [27] have consistently revealed that high hardness and strength properties achieved in high nitrogen martensitic stainless steel by tempering below 500 °C is attributable to solution strengthening effect of N and precipitation strengthening of extremely fine coherent nitride precipitates which can not be resolved using TEM, and subsequent sharp decrease in hardness and strength with increasing tempering temperature is caused by coarsening of precipitates and loss of coherence. This can well explain the slight influence of N addition on strength properties of 16Cr–5Ni–1Mo martensitic stainless steel upon tempering above 550 °C in present study. It is interesting to note that the yield strength of HN steels is lower than that of LN steel, and elongation of HN steels is higher than that of LN steel after tempering at temperatures of 650 °C and 700 °C. This is attributed to the retardation effect of N on formation of reversed austenite. The volume fraction of retained austenite in HN steels is higher than that in LN steel after tempering at these temperatures.
In summary, N as strong austenite stabilizing element can be used to partially substitute Ni to avoid d ferrite in 16Cr–5Ni–1Mo martensitic stainless steel. N alloyed super martensitic stainless steel exhibit improved strength properties and adequate toughness after normalizing treatment. To achieve excellent combined mechanical properties for high nitrogen martensitic stainless steel, dissolution of N in matrix or extremely fine coherent precipitation of nitrides and adequate amount of retained austenite are essential. Conventional heat treatment procedure of tempering above 550 °C is no more suitable for nitrogen alloyed super martensitic stainless, as N is found to retard the kinetics of reversed austenite formation and to promote precipitation of Cr2N at martensitic inter-lath boundaries after tempering treatment. Moreover, it is suspected that significant depletion of Cr would occur in vicinity of Cr2N at lath boundaries, which will be detrimental to the corrosion resistance. A novel heat treatment approach for high nitrogen super martensitic stainless steel will be exploited. 4. Conclusions (1) N as strong austenite stabilizing element can be used to partially replace Ni to effectively suppress the occurrence of d ferrite in martensitic stainless steel. The strength properties of the steel as normalized is enhanced by 0.1% N addition without decrease in ductility and toughness due to combined effects of solution strengthening of N, suppression d ferrite and retained austenite. (2) Precipitation of rod-like Cr2N along martensite lath boundaries and within laths of martensitic stainless steel with N after tempering above 550 °C contributes to slight increase in strength and much decrease in toughness compared to the steel without N. The effect of increasing tempering temperature is to increase the volume fraction of retained austenite originated from reversed austenite, which is very effective in restoring the ductility and toughness of the steels. N is found to retards the kinetics of formation of reversed austenite, thereby retarding the softening of the steel during tempering. (3) About 0.1% N addition to 16%Cr5%Ni1%Mo martensitic stainless steel offers good strength properties and decreased ductility and toughness after tempering due to the severe precipitation of Cr2N. Tempering treatment above 550 °C is undesirable for high N alloyed 16Cr5Ni1Mo martensitic stainless steel to achieve good combined mechanical properties.
Acknowledgment We acknowledge with grateful thanks for the award of a scholarship by China Scholarship Council to Xiaoping Ma to do research at McMaster University. References [1] Linne CP, Blanchard F, Guntz GC, Orlans-Joliet BJ. Corrosion performances of modified 13Cr for OCTG in oil and gas environments. Corrosion 1997;97:28. [2] Wang P, Lu SP, Xiao NM, Li DZ, Li YY. Effect of delta ferrite on impact properties of low carbon 13Cr–4Ni martensitic stainless steel. Mater Sci Eng A 2010;527:3210–6. [3] Carrouge D, Bhadeshia HKDH, Woollin P. Effect of d ferrite on impact properties of super martensitic stainless steel heat affected zones. Sci Technol Weld Join 2004;9:377–89. [4] Takano K, Sakakibara M, Matsui T, Takaki S. Effect of alloying elements and microstructure on corrosion resistance and hardness of martensitic stainless steel. J Iron Steel Inst Jpn 2008;86:123–30. [5] Kim SH, Moon HK, Kang T, Lee CS. Dissolution kinetics of delta ferrite in AISI 304 stainless steel produced by strip casting process. Mater Sci Eng A 2003;356:390–8.
X.P. Ma et al. / Materials and Design 34 (2012) 74–81 [6] Bashu SA, Singh K, Rawat MS. Effect of heat treatment on mechanical properties and fracture behavior of 12CrMoV steel. Mater Sci Eng A 1990;127:7–15. [7] Olson DL. Prediction of austenitic weld metal microstructure and properties. Welding J 1985;64:281–95. [8] Hamano S, Shimizu T, Noda T. Properties of low carbon high nitrogen martensitic stainless steels. Mater Sci Forum 2007;539–543:4975–80. [9] Bahrami F, Hendry A. Microstructure and mechanical behavior of high nitrogen alloyed martensitic stainless steel. Mater Sci Technol 1995;11:488–97. [10] Toro A, Misiolek WZ, Tschiptschin AP. Correlations between microstructure and surface properties in a high nitrogen martensitic stainless steel. Acta Mater 2003;51:3363–74. [11] Gocmen A, Ernst P, Holmes P. Principles of alloy design in high nitrogen 12% chromium steels. Mater Sci Forum 1999;318–320:215–26. [12] Ojima M, Ohnuma M, Suzuki J, Ueta S, Narita S, Shimizu T, et al. Origin of the enhanced hardness of a tempered high nitrogen martensitic stainless steel. Scr Mater 2008;59:313–6. [13] Horovitz MB, Neto FB, Garbogini A, Tschiptschin AP. Nitrogen bearing martensitic stainless steel: microstructure and properties. ISIJ Int 1996;36:840–5. [14] Calliari I, Zanesco M, Dabala M, Brunelli K, Ramous E. Investigation of microstructure and properties of a Ni–Mo martensitic stainless steel. Mater Des 2008;29:246–50. [15] Ma X, Wang L, Liu C. Smelting of high nitrogen martensitic stainless steels by vacuum induction furnace under near normal protective atmosphere. J Mater Metall 2009;8:168–71. [16] GB/T 28-2002. Metallic materials-tensile testing at ambient temperature; 2002.
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[17] ASTM A370. Stand test methods and definitions for mechanical testing of steel products; 2011. [18] Leem DS, Leeb YD, Jun JH, Choi CS. Amount of retained austenite at room temperature after reverse transformation of martensite to austenite in an Fe– 13%Cr–7%Ni–3%Si martensitic stainless steel. Scr Mater 2001;45:767–72. [19] Lee SJ, Park JS, Lee YK. Effect of austenite grain size on the transformation kinetics of upper and lower bainite in low-alloy steel. Scr Mater 2008;59:87–90. [20] Kaluba WJ, Kaluba T, Taillard R. The austenitizing behavior of high-nitrogen martensitic stainless steels. Scr Mater 1999;41:1289–93. [21] Park ES, Yoo DK, Sung JH, Kang CY, Lee J, Sung JH, et al. Met Mater Int 2004;6:521–5. [22] Al Dawood M, El Mahallawi IS, Abd El Azim ME, El Koussy MR. Thermal aging of 16Cr–5Ni–1Mo stainless steel. Part 1: Microstructural analysis. Mater Sci Technol 2004;20:363–9. [23] Qin B, Wang ZY, Sun QS. Effect of tempering temperature on properties of 00Cr16Ni5Mo stainless steel. Mater Charact 2008;59:1096–100. [24] Bilmes PD, Solari M, Llorente CL. Characteristics and effects of austenite resulting from tempering of 13Cr–NiMo martensitic steel weld metals. Mater Charact 2001;46:285–96. [25] Keller C, Margulies MM, Hadjem-Hamouche Z, Guillot I. Influence of the temperature on the tensile behaviour of a modified 9Cr–1Mo T91 martensitic steel. Mater Sci Eng A 2010;527:6758–64. [26] Berns H, Krafft F. Properties of high nitrogen steels. In: Proc 2nd int conf high nitrogen steels; 1999. p. 338–42. [27] Ma XP, Wang LJ, Qin B, Liu CM, Subramanian SV. Microstructure evolution and precipitation behavior of a high nitrogen 15Cr martensitic stainless steel during tempering. Adv Mater Res 2011;311–313:910–5.
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Effect of N on microstructure and mechanical properties of 16Cr5Ni1Mo martensitic stainless steel X.P. Ma a, L.J. Wang a, B. Qin b, C.M. Liu a,⇑, S.V. Subramanian c a
Key Laboratory for Anisotropy and Texture of Materials, Northeastern University, Shenyang 110004, China Research Institute for Stainless Steel, R&D Center, Baoshan Iron & Steel Co., Ltd., Shanghai 201900, China c Department of Materials Science and Engineering, McMaster University, Hamilton, Canada L8S-4L7 b
a r t i c l e
i n f o
Article history: Received 26 May 2011 Accepted 27 July 2011 Available online 3 August 2011 Keywords: A. Ferrous metals and alloys E. Mechanical F. Microstructure
a b s t r a c t The effect of adding N (about 0.1% in mass) to low carbon 16Cr5Ni1Mo martensitic stainless steel by adding nitride ferroalloy under argon protective atmospheric of near normal pressure on phase transformation and microstructure of steels subjected to normalizing and tempering was investigated by using dilatometer, laser scanning confocal microscopy, transmission electron microscopy and X-ray diffraction, and its consequence on mechanical properties was evaluated by tensile and impact tests. N can effectively avoid the occurrence of d ferrite during solidification. N retards the kinetics of the formation of reversed austenite and promotes the occurrence of Cr2N during tempering above 550 °C. The strength properties of the steel as normalized is enhanced by N alloying without much decrease in elongation and toughness due to the combined effects of N solution strengthening, suppression of d ferrite and retained austenite. The slight influence of N addition on mechanical properties of martensitic stainless steel after tempering above 550 °C is attributed to the precipitation of coarse rod-like Cr2N and loss of coherence. While the increase in volume fraction of retained austenite originated from reversed austenite with tempering temperature contributes to the decrease in strength properties and increase in toughness and elongation, the retransformation of reversed austenite into martensite affects these properties inversely. Ó 2011 Elsevier Ltd. All rights reserved.
1. Introduction Super low carbon martensitic stainless steels called super martensitic stainless steel (SMSS) are currently produced for the applications of oil country tubular goods (OCTG) [1]. They are based on the Fe–Cr–Ni–Mo system with 13–16 mass% Cr, 4–6 mass% Ni, 0.5– 2.5 mass% Mo. The basic design concept is the increase in effective Cr content by reducing C to improve corrosion resistance and weldability, the addition of Ni to maintain martensitic microstructure without d ferrite upon cooling and to achieve reversed austenite during tempering, and the addition of Mo to improve the resistance to localized corrosion and sulfide stress cracking (SSC). Martensitic microstructure can be achieved by air cooling from austenite single field due to their adequate hardenability. Retained austenite originated from reversed austenite during tempering is very effective in contributing to the high toughness and ductility of SMSS, which is another advantage of SMSS besides corrosion resistance. There has been a commercial 16Cr5Ni1Mo martensitic stainless steel. Higher chromium content of 16% in the commercial steel is designed to slow down the overall corrosion rate of the steel. How⇑ Corresponding author. Tel.: +86 24 83687554; fax: +86 24 23890920. E-mail address: [email protected] (C.M. Liu). 0261-3069/$ - see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2011.07.064
ever, Cr is a ferrite former and an excessive Cr content causes d ferrite to be generated, which not only impairs the toughness properties of the steel but also can generate SCC originated from the inter-phase between the martensite and ferrite [2–4]. Once the d ferrite forms, it is difficult to remove it by conventional heat treatments (except for high temperature diffusion annealing [5,6], which causes burning loss and severe distortion of castings). Thereby, it is necessary to make alloying additions so that austenite phase is stabilized in order for avoiding d ferrite and achievement of more retained austenite along martensitic boundaries and within laths during tempering to get toughness. Ni is the common alloying element used to stabilize the austenite or expand the c phase field, but more Ni application will add on cost. According to Delong and Reid diagram [7], the nickel equivalent number has to be increased by three to balance chromium equivalent number of 17 in commercial 16Cr–5Ni–1Mo martensitic stainless steel to avoid the formation of d ferrite. N is a cost effective austenite stabilizer, which is about 25–30 times more effective than Ni [7], and about 0.1 wt% addition of N is expected to effectively avoid the occurrence of d ferrite in 16Cr–5Ni–1Mo martensitic stainless steel. However, the low solubility of N in martensite make the additions of N in martensitic stainless steel difficult, and several special routes for production of high nitrogen martensitic stainless
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to the thickness of 100 lm and electropolished with a double jet Tenupol Struers machine at 40 °C, 15 V in an electrolyte containing 27 ml perchloric acid and 273 ml alcohol. Carbon extraction replica were prepared on sample surface well-polished and lightly etched using electrolyte containing 2 ml HCl + 4 ml HNO3 + 12 ml alcohol. The precipitate morphology, size and dispersion were examined using conventional transmission electron microscopy (TEM-Philips CM12 120 kV). The micro chemical compositions of precipitates were analyzed using Energy Dispersive Spectrometer (EDS) attached to TEM. The volume fraction of the retained austenite was measured at room temperature by X-ray diffraction using Cu Ka radiation on tempered specimens. The evaluation of volume fraction was determined by measuring the integrated intensities of (1 1 1)c and (1 1 0)a peaks of X-ray pattern, using the procedure outlined by Leem et al. [18].
steels has been used, such as pressurized-electro-slag-remelting [8], powder metallurgy and thermo-chemical treatments [9]. Meanwhile, a lot of works has been done to study the effect of introduction of N on microstructure and mechanical properties of martensitic stainless steel. In general, these works mainly aimed to combine elevated hardness and wear resistance with good corrosion resistance by substitution of C by N through more homogeneous precipitation of finely dispersed Cr nitride in Fe–Cr–N and Fe–Cr–Mo–N alloys [8–13]. Calliari et al. [14] has also studied the microstructure and properties of a martensitic stainless steel alloyed with 0.08 mass% N, but no specific effect of N was concerned. The present study tentatively use N to partially replace Ni. The effect of introduction of N into commercial super low carbon 16Cr– 5Ni–1Mo martensitic stainless steel under argon protective atmosphere on microstructure evolution during heat treatment, and its consequence on mechanical properties including tensile and Chary impact toughness were investigated to generate base line information for high nitrogen Fe–Cr–Ni–Mo martensitic stainless steel.
3. Results and discussion 3.1. Microstructure characterization
2. Experimental procedure Fig. 1 shows the LSCM images of LN and HN2 steels as normalized disclosing martensitic microstructure achieved by air cooling the steels from austenitizing temperature of 1100 °C. Prior austenite grains were subdivided into packets of martensite laths. About 6.5% d ferrite occurs along prior austenite grain boundaries of LN steel due to its higher Cr equivalent. Compared with LN steel, the amount of d ferrite is significantly decreased to 1.2% and 0.7% in HN1 and HN2 steels, respectively by introduction of 0.12 wt% and 0.096 wt% N to LN steel due to strong austenite stabilizing effect of effect N. This is in agreement with original design concept. N can be used to replace Ni to avoid d ferrite effectively. Careful TEM observation on as normalized samples with high nitrogen did not reveal the presence of N rich precipitates. It is supposed that N was entirely dissolved into austenite during solution treatment at 1100 °C and did not precipitate during air cooling. Fig. 2 shows the TEM image of detailed microstructure of LN steel as normalized showing highly dislocated martensitic laths of 0.2–0.3 lm thick and d ferrite. N as very strong austenite stabilizer may increase the volume fraction of retained austenite after normalizing, which will increase the toughness of the steel, but decrease the strength. However, X-ray diffraction test could not prove this point due to low volume fraction of retained austenite in as normalized samples. The comparison of X-ray diffraction pattern of three steel samples as normalized did not found apparent influence of N on lattice parameter, and the structure of N alloyed martensite remains body-centered cubic, other than tetragonal. Fig. 3 shows the austenite transformation volume fraction as a function of heating temperature calculated from the continuous heating curve of dilation diagrams of as-normalized LN and HN2 steel samples at heating rate of 3 and 60 °C/min, using the procedure outlined by Lee et al. [19]. As can be seen, the curves of the same steel at higher heating rate are shifted towards higher temperatures, which means that transformation involved in the austenitization process is heating rate sensitive. At lower heating rate, the phase transformation temperature will get closer to equilibrium point. The curves are shifted towards high temperature for HN2 steel compared with LN steel, demonstrating that addition of
High nitrogen steel samples were made in laboratory in the form of ingots by melting in a 100 kg vacuum induction furnace under argon protective atmosphere of near normal pressure through adding nitride ferroalloy (nitrogen of 5%) and electroslag remelting under atmospheric pressure [15]. The chemical compositions of three tested steels with and without N are summarized in Table 1. The content of N was determined using the spark emission ARL 4460 spectrometer. While commercial LN steel is the control steel without N, HN1 and HN2 are indicated as high nitrogen stainless steel. The ingots were hot rolled at 1200 °C into plates with 12 mm thickness. Normalizing was carried out at 1100 °C for 0.5 h, followed by tempering for 2 h at temperatures from 550 °C to 700 °C. After tempering, the samples were quenched in oil. The heating rate of tempering treatment is about 60 °C/min. The dilatometer technique was used to study the thermal expansion of the steels during continuous heating treatments, which were carried out by heating as-normalized samples from room temperature at different rate of 3 °C/min and 60 °C/min to the austenitizing temperature of 1100 °C. The mechanical properties were evaluated by means of tensile and impact tests. The tensile tests were performed on a tensile testing machine CMT5105 with a extensometer at strain rate of 0.001 s 1. Cylindrical test specimens with a uniform gauge cross-section of U 5 25 mm and machined parallel to the plate rolling direction were used [16]. The impact tests were performed on full size Charpy V notch impact specimens (10 10 55 mm) at room temperature [17]. All mechanical properties value reported in the results are the average of three tests. Metallographic specimens were cut from broken impact samples. The microstructures of samples were investigated by using Laser scanning confocal microscopy (LSCM) and transmission electron microscopy (TEM). The fraction of d ferrite was estimated by quantitative image analysis software at 200 magnification according to the contrast between d ferrite and martensite matrix, and the final value reported is the average of five measurements. Thin foils for TEM analysis were prepared by mechanically grinding the samples with initial thickness of 400 lm
Table 1 Chemical compositions of the tested steels (in mass%). Steel grade
C
Si
Mn
P
S
Cr
Ni
Mo
N
Fe
LN HN1 HN2
0.03 0.017 0.015
0.32 0.32 0.32
0.65 0.78 0.8
0.011 0.011 0.011
0.011 0.007 0.008
16.37 16.32 16.26
5.12 5.03 5.01
0.83 0.98 1.03
0.0079 0.096 0.12
Bal. Bal. Bal.
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Fig. 1. LSCM images of (a) LN and (b) HN2 steels as normalized showing martensitic microstructure and d ferrite distributed along prior austenite grain boundaries.
Fig. 2. TEM image of LN steel as normalized from austenitizing temperature of 1100 °C showing d ferrite and highly dislocated martensitic laths of 0.2–0.3 lm in thickness.
Fig. 3. The variation of austenite volume fraction with heating temperature calculated from continuous heating dilation diagram of two SMSSs with heating rate of 3 and 60 °C/min.
N retards the austenite transformation. This is contrary with the original purpose of design of high nitrogen martensitic stainless steel. N as strong austenite stabilizing element was expected to promote the formation of reversed austenite. According to the
work by Kaluba et al. [20] on martensitic stainless steel with nitrogen, the process of austenitizing of martensite during heat treatment is mainly affected by diffusion of interstitial element C and N. It is proposed in present study that the kinetics of austenite transformation is restrained due to the additional diffusion of N. As determined in Fig. 3, although the Ac1 temperature is further above 550 °C up to 625 °C and 850 °C at heating rate of 3 °C/min and 60 °C/min, respectively, the reverse transformation from martensite to austenite would occur if the samples were maintained above 550 °C in present study. In agreement with many studies carried out on microstructure evolution of martensitic stainless steel during tempering [21,22], reversed austenite obtained by tempering the martensite is relatively stable upon cooling to room temperature and can be partially retained at room temperature. Previously published work by Qin et al. [23] on LN steel has shown that the volume fraction of retained austenite in samples tempered at various temperatures for 2 h increases with increasing tempering temperature and peaks after tempering at 600 °C, above which the volume fraction of retained austenite decreases due to the retransformation of reversed austenite to martensite owing to lowered stability of reversed austenite after tempering at higher temperature. It has been reported that the stability of reversed austenite during cooling after tempering relates to its chemical compositions as well as tempering temperatures [24]. The concentration of alloy elements in reversed austenite during tempering makes it stable due to the decreased MS. With higher tempering temperature, the quenched-in vacancies are expected to increase, which would lower the stability of reversed austenite to be retained upon cooling. Most of reversed austenite can be retained after tempering at lower temperature due to its higher stability. As the temperature is increased further, the reversed austenite volume fraction increases rapidly so that it is only slightly enriched in alloy content and largely retransformed into ‘‘fresh’’ martensite on cooling. Fig. 4a shows the results of X-ray diffraction carried out using Cu Ka radiation on samples of three steels tempered at 600 °C for 6 h. According to the calculation procedure outline by Leem [18], the volume fraction of retained austenite is about 14.4%, 9.3%, 7.2% in LN, HN1 and HN2 steel, respectively. Due to the retardation effect of N on formation of reversed austenite, the tempering temperature for certain volume fraction of reversed austenite, at which the concentration of alloying elements in reversed austenite decreases and composition of reversed austenite is close to average contents of the steel, is higher for HN steels than for the LN steel, thus the retransformation of reversed austenite to martensite would occur at different extent after tempering at same temperature, more in LN steel. As shown in Fig. 4b, while the peaks of retained austenite are present in X-ray diffraction pattern of HN steel tempered at 650 °C for 2 h, the diffraction pattern on LN steel
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Fig. 4. X-ray diffraction results obtained by Cu Ka radiation on samples tempered at (a) 600 °C for 6 h and (b) 650 °C for 2 h.
sample tempered at same temperature for 2 h shows no retained austenite peak, demonstrating that the volume fraction of austenite retained in HN2 steel is higher than that in LN steel after tempering at 650 °C. Fig. 5a and b exhibit TEM bright field and dark field pictures of retained austenite of 100–200 nm in thickness in HNl steel after tempering at temperature of 600 °C. Fig. 5c and d show TEM morphologies of retained austenite in HN1 steel after tempering at 700 °C using bright field and dark field, respectively. The retained austenite film is found with thickness of about 50 nm. Due to the limitation of equipment, the segregation of N between retained austenite and matrix can not be exactly detected.
Previous studies [2,25] have shown that the precipitates in low carbon martensitic stainless steel during tempering are mainly M23C6, where M represents Cr, Mo and Fe. It is reported that the addition of nitrogen delays carbide precipitation and results in disappearance of carbide phase in high nitrogen martensitic stainless steel [26]. Previous work [27] by the authors on microstructures evolution of 15Cr martensitic stainless steel with 0.16 wt% N during tempering has shown that while tempering below 500 °C does not show Cr2N precipitation under TEM observation, tempering above 500 °C results in severe precipitation of Cr2N along martensitic inter-lath boundaries and within matrix. Although the solubility limit of N in martensitic stainless steel is unknown, recent study
Fig. 5. (a) and (b) Bright field and dark field TEM images of HN1 sample tempered at 600 °C showing retained austenite with thickness of 100–200 nm distributed along lath boundaries and within laths; (c) and (d) bright field and dark field TEM images of HN1 sample tempered at 700 °C showing martensite laths and retained austenite with thickness of 50 nm.
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found that the thermodynamic potential for precipitation of Cr2N even occur in 13Cr–5Ni–1Mo martensitic stainless steel with only 0.04 wt% N addition. Fig. 6 shows the TEM images of morphologies, dispersion of precipitates observed in HN1 steel tempered at various temperatures. Fig. 6a shows the TEM pictures of rod-like pre-
cipitates occurring at martensitic lath boundaries and within the lath of HN1 steel tempered at 550 °C. Fig. 6b and c show the TEM pictures of precipitates within martensite lath and at lath boundaries in HN1 steel tempered at 600 °C, respectively. The rod-like precipitates within matrix is around 10 nm in diameter and
Fig. 6. TEM pictures of HN1 steel tempered at: (a) 550 °C showing precipitates along lath boundary and within the lath; (b) and (c) 600 °C showing precipitates distributed randomly within matrix and along lath boundaries; (d) and (e) TEM pictures of carbon extraction replica of HN1 steel tempered at 600 °C showing precipitates well aligned in arrays and randomly distributed, respectively; (f) TEM images of thin foil of HN1 steel tempered at 700 °C showing boundary precipitates of 15 nm in diameter and 30–80 nm in length; (g) ED spectrum analysis of precipitates exhibited in (d) and (e) showing X-ray signals characteristic of Cr and small amount of Mo; (h) SADP of precipitates shown in (e).
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50 nm in length, and precipitates at lath boundaries is continuous. Due to low contrast between precipitates and matrix, and the interference from matrix, it is difficult to observe and identify the precipitates using thin foil. Fig. 6d and e show TEM pictures of carbon extraction replica of HN1 steel tempered at 600 °C showing precipitates well aligned in arrays with inter-spacing comparable with inter-lath spacing of martensitic lath shown in Fig. 6c and precipitates randomly dispersed, respectively. It is considered that the precipitates well aligned were nucleated at boundaries and the precipitates randomly dispersed were nucleated within martensitic laths. The ED spectrum of all precipitates shows X-ray signals characteristic of Cr, Mo and Fe, as shown in Fig. 6g, where Cu arises from copper grid. From the selected area diffraction pattern (SADP) shown in Fig. 6h, these precipitates are identified with hcp structure with parameters of a = 0.480 nm, c = 0.447 nm. According to the crystalline structure and lattice parameter as well as the EDS analysis, these precipitates are identified as Cr2N. It is suspected that significant depletion of Cr in vicinity of Cr rich precipitates at lath boundaries would occur due to high diffusion rate of substantial element at boundaries, which will deteriorate corrosion resistance. Fig. 6f shows TEM picture of HN1 steel tempered at 700 °C showing Cr2N of 15 nm in diameter and 30–80 nm in length. With increasing tempering temperature, Cr2N grew slightly. Actually, the precipitation of Cr2N in this temperature range is undesirable. However, in order to get more retained austenite to restore the ductility and toughness, it is necessary to tempering the steel above 550 °C at which Cr2N would precipitate simultaneously. 3.2. Mechanical properties Fig. 7 shows the tensile properties, Charpy impact toughness of three tested steels at room temperature after normalizing and tem-
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pering at temperatures varied from 550 °C to 700 °C for 2 h. The yield strength is 775, 795, 815 MPa, and the ultimate tensile strength is 1005, 1055 and 1115 MPa for LN, HN1 and HN2 steels, respectively in normalized condition. It is noted that while the yield strength and ultimate tensile strength of the steels as normalized is enhanced to some extent by N addition in two HN steels, the elongation and Charpy impact toughness is slightly improved for HN1 and decreased for HN2 steel. Bahrami and Hendry [9] reported that the strength increases linearly with increasing nitrogen content in 10% Cr martensitic stainless steel in solution treated condition. This is inconsistent with the results in present study. Combined with comparison of microstructure of the steels with and without N presented above, the difference in mechanical properties between steels with and without N after normalization is not only attributed to the solution strengthening effect of N, but also suppression of d ferrite and possible increase in volume fraction of retained austenite by N. Wang et al. [2] found d ferrite is soft owing to its relative low C content and dislocation density in low carbon martensitic stainless steel. Suppression of d ferrite plus solution strengthening effect of N is supposed to significantly increase the strength of the N alloyed martensitic stainless steel. However, the increment in strength caused by N addition is much lower than that reported in literatures. In addition, Carrouge et al. [3] and Wang consistently proposed that the presence of d ferrite in low C martensitic stainless steel does not change the upper and lower shelf energy apparently, while lowers the impact energy remarkably in the transition temperature range and raises the ductile to brittle transition temperature. As the temperature of 15 °C employed in testing the Charpy impact toughness in present study is above the transition temperature range reported in literatures, the suppression of d ferrite is not considered to be the main factor affecting the tensile properties and impact toughness at room tem-
Fig. 7. Mechanical properties of the three studied steels as a function of tempering temperature at room temperature. (a) Yield strength (Rp0.2). (b) Ultimate tensile strength (Rm). (c) Elongation (A%) and (d) Charpy impact toughness (Akv).
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perature. Combining above effects of N on tensile strength properties and Chary toughness, it is deducted that the increase in retained austenite volume fraction in N alloyed martensitic stainless steel in as normalized condition significantly contributes to the improved toughness and slightly increased strength of N alloyed steels in comparison with the control alloy without N. While tempering above 550 °C promotes the precipitation of Cr2N within martensitic laths and at martensitic lath boundaries, strengthening the alloy at cost of elongation and toughness, retained austenite is formed along lath boundaries and within laths simultaneously, restoring ductility and toughness. The mechanical properties of the tested steels after tempering depend on the balance among softening effect of retained austenite, recovery effect of martensite matrix, and the hardening effect of precipitates and the retransformation of reversed austenite to martensite. As shown in Fig. 7a, the yield strength of three tested steels is increased by tempering at 550 °C in relation to as normalized condition due to the precipitation of Cr carbide in LN steel and Cr nitride in HN steels. The ductility and toughness are effectively improved by tempering treatment, owing to the recovery of martensite matrix and formation of small amount of retained austenite during tempering, as shown in Fig. 7c and d. As tempering temperature is increased, the yield strength and ultimate tensile strength first decrease to minimum values after tempering at 625 °C, 650 °C and 600 °C, 625 °C for LN and HN steels, respectively, and then revert to increase, while the elongation and toughness change contrarily, first increasing to maximum values at 600 °C, 625 °C and 600 °C, 650 °C, respectively and then revert to decrease. Retained austenite originated from reversed austenite is reported largely to be very effective in increasing toughness and elongation. The increase in volume fraction of retained austenite apparently improves ductility and toughness, and its decrease lowers both ductility and toughness. This is also can be confirmed in present study. Combined with the measurement of retained austenite volume fraction in LN steel presented previously, the temperature at which the elongation and toughness of LN steel are highest is in agreement with the temperature at which the highest volume fraction of retained austenite has been registered. Deducing from this point, the increase in strength properties and the decrease in ductility and toughness of HN steels after tempering at higher temperature are mainly attributed to the retransformation of reversed austenite to martensite. On the other hand, the retardation of N on formation of reversed austenite would retard the softening of martensite during tempering. Comparing all mechanical properties of three tested steels tempered in the temperature range from 550 °C to 700 °C, the strength properties is slightly enhanced, while the Charpy toughness of steel is severely decreased by N alloying due to the precipitation of Cr2N at lath boundaries and within matrix. Previous work by other researchers [9,12] and the author [27] have consistently revealed that high hardness and strength properties achieved in high nitrogen martensitic stainless steel by tempering below 500 °C is attributable to solution strengthening effect of N and precipitation strengthening of extremely fine coherent nitride precipitates which can not be resolved using TEM, and subsequent sharp decrease in hardness and strength with increasing tempering temperature is caused by coarsening of precipitates and loss of coherence. This can well explain the slight influence of N addition on strength properties of 16Cr–5Ni–1Mo martensitic stainless steel upon tempering above 550 °C in present study. It is interesting to note that the yield strength of HN steels is lower than that of LN steel, and elongation of HN steels is higher than that of LN steel after tempering at temperatures of 650 °C and 700 °C. This is attributed to the retardation effect of N on formation of reversed austenite. The volume fraction of retained austenite in HN steels is higher than that in LN steel after tempering at these temperatures.
In summary, N as strong austenite stabilizing element can be used to partially substitute Ni to avoid d ferrite in 16Cr–5Ni–1Mo martensitic stainless steel. N alloyed super martensitic stainless steel exhibit improved strength properties and adequate toughness after normalizing treatment. To achieve excellent combined mechanical properties for high nitrogen martensitic stainless steel, dissolution of N in matrix or extremely fine coherent precipitation of nitrides and adequate amount of retained austenite are essential. Conventional heat treatment procedure of tempering above 550 °C is no more suitable for nitrogen alloyed super martensitic stainless, as N is found to retard the kinetics of reversed austenite formation and to promote precipitation of Cr2N at martensitic inter-lath boundaries after tempering treatment. Moreover, it is suspected that significant depletion of Cr would occur in vicinity of Cr2N at lath boundaries, which will be detrimental to the corrosion resistance. A novel heat treatment approach for high nitrogen super martensitic stainless steel will be exploited. 4. Conclusions (1) N as strong austenite stabilizing element can be used to partially replace Ni to effectively suppress the occurrence of d ferrite in martensitic stainless steel. The strength properties of the steel as normalized is enhanced by 0.1% N addition without decrease in ductility and toughness due to combined effects of solution strengthening of N, suppression d ferrite and retained austenite. (2) Precipitation of rod-like Cr2N along martensite lath boundaries and within laths of martensitic stainless steel with N after tempering above 550 °C contributes to slight increase in strength and much decrease in toughness compared to the steel without N. The effect of increasing tempering temperature is to increase the volume fraction of retained austenite originated from reversed austenite, which is very effective in restoring the ductility and toughness of the steels. N is found to retards the kinetics of formation of reversed austenite, thereby retarding the softening of the steel during tempering. (3) About 0.1% N addition to 16%Cr5%Ni1%Mo martensitic stainless steel offers good strength properties and decreased ductility and toughness after tempering due to the severe precipitation of Cr2N. Tempering treatment above 550 °C is undesirable for high N alloyed 16Cr5Ni1Mo martensitic stainless steel to achieve good combined mechanical properties.
Acknowledgment We acknowledge with grateful thanks for the award of a scholarship by China Scholarship Council to Xiaoping Ma to do research at McMaster University. References [1] Linne CP, Blanchard F, Guntz GC, Orlans-Joliet BJ. Corrosion performances of modified 13Cr for OCTG in oil and gas environments. Corrosion 1997;97:28. [2] Wang P, Lu SP, Xiao NM, Li DZ, Li YY. Effect of delta ferrite on impact properties of low carbon 13Cr–4Ni martensitic stainless steel. Mater Sci Eng A 2010;527:3210–6. [3] Carrouge D, Bhadeshia HKDH, Woollin P. Effect of d ferrite on impact properties of super martensitic stainless steel heat affected zones. Sci Technol Weld Join 2004;9:377–89. [4] Takano K, Sakakibara M, Matsui T, Takaki S. Effect of alloying elements and microstructure on corrosion resistance and hardness of martensitic stainless steel. J Iron Steel Inst Jpn 2008;86:123–30. [5] Kim SH, Moon HK, Kang T, Lee CS. Dissolution kinetics of delta ferrite in AISI 304 stainless steel produced by strip casting process. Mater Sci Eng A 2003;356:390–8.
X.P. Ma et al. / Materials and Design 34 (2012) 74–81 [6] Bashu SA, Singh K, Rawat MS. Effect of heat treatment on mechanical properties and fracture behavior of 12CrMoV steel. Mater Sci Eng A 1990;127:7–15. [7] Olson DL. Prediction of austenitic weld metal microstructure and properties. Welding J 1985;64:281–95. [8] Hamano S, Shimizu T, Noda T. Properties of low carbon high nitrogen martensitic stainless steels. Mater Sci Forum 2007;539–543:4975–80. [9] Bahrami F, Hendry A. Microstructure and mechanical behavior of high nitrogen alloyed martensitic stainless steel. Mater Sci Technol 1995;11:488–97. [10] Toro A, Misiolek WZ, Tschiptschin AP. Correlations between microstructure and surface properties in a high nitrogen martensitic stainless steel. Acta Mater 2003;51:3363–74. [11] Gocmen A, Ernst P, Holmes P. Principles of alloy design in high nitrogen 12% chromium steels. Mater Sci Forum 1999;318–320:215–26. [12] Ojima M, Ohnuma M, Suzuki J, Ueta S, Narita S, Shimizu T, et al. Origin of the enhanced hardness of a tempered high nitrogen martensitic stainless steel. Scr Mater 2008;59:313–6. [13] Horovitz MB, Neto FB, Garbogini A, Tschiptschin AP. Nitrogen bearing martensitic stainless steel: microstructure and properties. ISIJ Int 1996;36:840–5. [14] Calliari I, Zanesco M, Dabala M, Brunelli K, Ramous E. Investigation of microstructure and properties of a Ni–Mo martensitic stainless steel. Mater Des 2008;29:246–50. [15] Ma X, Wang L, Liu C. Smelting of high nitrogen martensitic stainless steels by vacuum induction furnace under near normal protective atmosphere. J Mater Metall 2009;8:168–71. [16] GB/T 28-2002. Metallic materials-tensile testing at ambient temperature; 2002.
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[17] ASTM A370. Stand test methods and definitions for mechanical testing of steel products; 2011. [18] Leem DS, Leeb YD, Jun JH, Choi CS. Amount of retained austenite at room temperature after reverse transformation of martensite to austenite in an Fe– 13%Cr–7%Ni–3%Si martensitic stainless steel. Scr Mater 2001;45:767–72. [19] Lee SJ, Park JS, Lee YK. Effect of austenite grain size on the transformation kinetics of upper and lower bainite in low-alloy steel. Scr Mater 2008;59:87–90. [20] Kaluba WJ, Kaluba T, Taillard R. The austenitizing behavior of high-nitrogen martensitic stainless steels. Scr Mater 1999;41:1289–93. [21] Park ES, Yoo DK, Sung JH, Kang CY, Lee J, Sung JH, et al. Met Mater Int 2004;6:521–5. [22] Al Dawood M, El Mahallawi IS, Abd El Azim ME, El Koussy MR. Thermal aging of 16Cr–5Ni–1Mo stainless steel. Part 1: Microstructural analysis. Mater Sci Technol 2004;20:363–9. [23] Qin B, Wang ZY, Sun QS. Effect of tempering temperature on properties of 00Cr16Ni5Mo stainless steel. Mater Charact 2008;59:1096–100. [24] Bilmes PD, Solari M, Llorente CL. Characteristics and effects of austenite resulting from tempering of 13Cr–NiMo martensitic steel weld metals. Mater Charact 2001;46:285–96. [25] Keller C, Margulies MM, Hadjem-Hamouche Z, Guillot I. Influence of the temperature on the tensile behaviour of a modified 9Cr–1Mo T91 martensitic steel. Mater Sci Eng A 2010;527:6758–64. [26] Berns H, Krafft F. Properties of high nitrogen steels. In: Proc 2nd int conf high nitrogen steels; 1999. p. 338–42. [27] Ma XP, Wang LJ, Qin B, Liu CM, Subramanian SV. Microstructure evolution and precipitation behavior of a high nitrogen 15Cr martensitic stainless steel during tempering. Adv Mater Res 2011;311–313:910–5.