- Email: [email protected]
Effect of solution treatment on cavitation erosion behavior of high-nitrogen austenitic stainless steel
Wear 424–425 (2019) 70–77
Contents lists available at ScienceDirect
Wear journal homepage: www.elsevier.com/locate/wear
Effect of solution treatment on cavitation erosion behavior of high-nitrogen austenitic stainless steel ⁎
T
⁎
Yanxin Qiaoa, , Jian Chena, , Huiling Zhoua, Yuxin Wanga, Qining Songb, Huabing Lic, Zhibin Zhengd a
School of Materials Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212003, China College of Mechanical and Electrical Engineering, Hohai University, Changzhou 213022, China c School of Metallurgy, Northeastern University, Shenyang 11004, China d Guangdong Key Laboratory for Technology and Application of Metal Toughening, Guangdong Institute of Materials and Processing, Guangzhou 510650, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: High-nitrogen austenitic stainless steel Solution treatment Cavitation erosion Elastic property
The effect of solution treatment on the cavitation erosion behavior of a Fe-18Cr-16Mn-2Mo-0.66 N high-nitrogen austenitic stainless steel (HNSS) exposed in distilled water up to 8 h was investigated using magnetostrictive cavitation, micro-hardness, mass loss measurements, scanning electron microscope (SEM) and X-ray diffraction (XRD). The results showed that the solution treatment slightly decreased the hardness of the HNSS but significantly increased its elasticity. The HNSS with an improved elastic property was shown to have a higher resistance to the absorption of energy produced by cavitation erosion, and hence a higher cavitation erosion resistance. The cavitation erosion damage of the HNSS in distilled water was found to be a deformation-controlled process which limited by the formation and development of plastic deformation during cavitation erosion tests.
1. Introduction Austenitic stainless steels have been widely used in various industrial fields because of their excellent mechanical and corrosion properties. However, the low yield strength limits their widespread application. In order to overcome this limitation, nitrogen is used as an additive in austenitic stainless steels [1–5]. High-nitrogen nickel-free stainless steels (HNSSs) have developed and gained immense attention for various industrial applications owing to their low cost [6], high strength and ductility [7], high work-hardening ability [8], high stressinduced martensite transformation ability [9], high corrosion resistance [10,11], and good biocompatibility [12]. Moreover, these alloys also show excellent cavitation erosion (CE) resistance [13–19], which depends on their microstructure and mechanical properties. Nitrogencontaining austenitic stainless steels also show high yield strength, ductility, and fracture toughness [1,3,20]. However, when exposed to temperatures higher than 900 °C for a short duration (i.e., 2 h), their mechanical properties would deteriorate due to the precipitation of Cr2N secondary phase and the formation of (Mn, Cr)-oxides in the HNSSs and nitrogen-containing austenitic stainless steels [21–24]. Solution treatment is widely used to homogenize metallic alloys, which can reduce the precipitation of the secondary phases, optimize ⁎
the microstructure, and improve the strength [25–27]. Although solution treatment significantly affects the microstructure, grain size, and mechanical properties of the alloys, the effect of solution treatment on the CE property of the metallic alloys still needs further investigations [28–30]. The solution treatment of UNS S30400 stainless steel decreased its CE resistance because of the increased grain size and ductility [28]. However, the effect of solution treatment on the CE behavior of S32760 duplex stainless steel was found be to minimal since the mechanical properties and grain size of the S32760 were not significantly affected by solution treatment [28]. Cheng et al. [31] investigated the CE behavior of the solution-treated and aged (aged after solution treatment) NiTi (50.8 at% Ni) and found that the CE resistance of the solution-treated sample was higher than that of the samples aged at temperatures in the range from 200 to 600 °C. The sensitization treatment of UNS S31803 duplex stainless steel at 475 °C deteriorated its CE resistance (~ 22% lower than that obtained after the solution annealing treatment) [32]. High-temperature (1300 °C) solution treatment of CrMnNi austenitic stainless steels led to a significant grain growth, which compromised their strength and facilitated the formation of martensite, and thereby deteriorated their CE properties [33]. Therefore, further studies on the CE process and mechanism are required to investigate the effect of heat treatment on the CE behavior of
Corresponding authors. E-mail addresses: [email protected] (Y. Qiao), [email protected] (J. Chen).
https://doi.org/10.1016/j.wear.2019.01.098 Received 2 November 2018; Received in revised form 29 January 2019; Accepted 29 January 2019 Available online 30 January 2019 0043-1648/ © 2019 Elsevier B.V. All rights reserved.
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
where hc is the contact depth at the peak load, P (h) is the corresponding load, ε is a correction index (0.75 for Berkovich indenter), and S is the slope of the unloading curve. The CE test was carried out in distilled water at 20 ± 1 °C. The dimensions of the specimen for the cavitation test and the details of the test were described elsewhere [13]. The CE test was repeated three times for each sample. The microstructures and eroded morphologies of the HNSS samples were observed using scanning electron microscopy (XL30-FEG ESEM). The crystal structures of the treated (solution treatment) and untreated HNSS samples were examined using a D8 Xray diffractionmeter (XRD) with a Cu Ka radiation at 40 kV and 35 mA. To investigate the effect of cavitation on work hardening of the samples, their micro-hardness was measured at a fixed interval using an MH-5 tester with a load of 100 g and a loading time of 15 s. These measurements were performed on the cross-sections of the specimens after CE for 1, 3, 5 and 8 h at a distance of 20 µm below the CE surface for five times. The average value of five measurements was used to represent the hardness of the eroded specimens.
Table 1 Mechanical parameters of the HNSS and S-HNSS.
HNSS S-HNSS
Rm (MPa)
Rp0.2 (MPa)
Hv100
920.4 ± 22.1 855.8 ± 16.3
570.6 ± 15.4 505.2 ± 14.3
299.5 ± 5.3 285.1 ± 4.5
metallic alloys such as the HNSSs. In this study, CE tests were carried out on a nickel-free HNSS (Fe18Cr-16Mn-2Mo-0.66 N) in distilled water using an ultrasonic vibration system. The effect of solution treatment on the CE properties and damage mechanism of this nickel-free HNSS were investigated. The erosion mechanisms were discussed. 2. Experimental details The chemical composition (wt%) of the Fe-18Cr-16Mn-2Mo-0.66 N HNSS used in this study was C 0.048%, Si 0.24%, Mn 15.96%, P 0.004%, S 0.017%, Cr 18.44%, Mo 2.23%, N 0.66%, and Fe balanced. The hot-forged ingots were cut into round bars with a diameter of 20 mm and were labeled as wrought HNSS. Some of the wrought HNSS samples were heated to 1150 °C for 1 h and were then quenched in water. These samples were named as S-HNSS. Tensile tests were carried out on an Instron-type testing machine with a tensile rate of 3 mm s-1 at 25 ± 1 °C. The mechanical properties of the HNSS and S-HNSS, including the surface hardness (Hv100), the tensile Rm and yield Rp0.2 strength are shown in Table 1. To characterize the elastic properties of the HNSS and S-HNSS, nanoindentation curves were acquired using a CSM NHT2 nanoindenter (Anton Paar) at a maximum load of 20 mN with a loading/unloading rate of 40 mN·min-1. In nanoindentation tests, the elastic behavior of a material can be analyzed by its depth recovery ratio (ηh) obtained from the load-displacement curve [34]. The value of ηh is calculated as follows:
ηh =
h max − hr h max
3. Results 3.1. Material characterization Fig. 1 shows the microstructure of the wrought HNSS and S-HNSS. The wrought HNSS was composed of an austenitic phase matrix (marked as 1 in Fig. 1(a)) with twins and homogeneously-distributed fine nitrides Cr2N particles (marked as 2 in Fig. 1(a)) formed during the solidification process [36]. The S-HNSS exhibited a single austenite phase (marked as 3 in Fig. 1(b)) with twins distributing inside the austenite grains. The grain size of the S-HNSS sample was larger (~25 µm) than that of the wrought HNSS sample (~8 µm). Fig. 2 shows the XRD patterns of the wrought HNSS and S-HNSS samples. Although both the wrought HNSS and S-HNSS were composed of a single γ phase with a face centered cubic (f.c.c.) structure, the characteristic peaks of the γ phase in the wrought HNSS were broader than those of the S-HNSS. According to the Scherrer equation [37], the grain size of a material is inversely proportional to the width of its peaks. Hence, the grain size of the wrought HNSS was smaller than that of the S-HNSS. This is consistent with the microscopic results shown in Fig. 1.
(1)
where hmax is the maximum penetration depth and hr is the residual depth after unloading. The hardness of stainless steels is defined as the ratio of the peak indentation load (Fmax) to the hardness impression projected area (Ac) and is calculated using Eqs. (2) and (3) as follows [35]:
F Fmax Hd = max = Ac 26.43hc2
hc = h max − ε
P (h) S
3.2. Nanoindentation test Fig. 3 shows the nanoindentation load-displacement curves (P-h) of the wrought HNSS and S-HNSS samples. The maximum indentation depth (hmax) at a load of 20 mN of the wrought HNSS sample (422.1 nm) was larger than that of the S-HNSS (377.2 nm). The other
(2) (3)
Fig. 1. The microstructure of the wrought HNSS (a) and S-HNSS (b). 71
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
and the mass loss rate of the wrought HNSS sample was larger than that of the S-HNSS sample. Both the wrought HNSS and S-HNSS samples demonstrated a plateau period and acceleration period [38]. Although the plateau periods for both the samples were observed during the initial 4 h, the mass loss rate of the wrought HNSS in this period was 3.4 times of that of the S-HNSS, indicating that the S-HNSS showed a high cavitation resistance in distilled water. Once the acceleration period of CE started, the erosion damage spread over the entire area of the specimen [39], leading to a rapid increase in the mass-loss rate. During the accumulation period, the slopes of the weight loss curves of the wrought HNSS and S-HNSS were 2.0 × 10-3 and 1.6 × 10-3 mg cm-2 h-1, respectively, indicating that the CE resistance of the HNSS was improved by the solution treatment.
3.3. Morphologies of the eroded surfaces Fig. 5 shows the SEM micrographs of the wrought HNSS and S-HNSS samples after erosion for 1 h. As shown in Fig. 5(a) and (c), the wrought HNSS sample experienced deformation in the presence of slip lines and deformed twins inside the grains, and surface undulations. In addition, the wrought HNSS suffered from material loss at the grain and twin boundaries. Some small holes with either circular or angular shapes were also observed on the eroded surface (black arrows in Fig. 5(c)). Taking into account the geometry of these holes, it is most likely that they were formed by the spalling of the Cr2N precipitation during the CE test. This suggests that the damage induced by CE was preferentially initiated at the grain boundaries, slip lines, and precipitates [40]. Although the S-HNSS sample exhibited plastic deformation, its material loss was minor, as shown in Fig. 5(b) and (d). This is consistent with the mass loss results shown in Fig. 4. This further demonstrated that solution treatment can improve the CE resistance of the HNSSs. The damaged surfaces of the wrought HNSS and S-HNSS samples after erosion for 3 h are shown in Fig. 6. With an increase in the CE duration, both the specimens experienced more pronounced deformation. The wrought HNSS exhibited obvious cavitation damage at the grain and twin boundaries, as shown in Fig. 6(a) and (c), which started to spread into the interior of the grains. Although some grains underwent intense plastic deformation with the establishment of multiple slips and initiation of micro-cracks inside the grains, some grains still remained intact and smooth as the original polished surface. This can be attributed to the different values of the resolved shear stress inside the grains [41,42]. At longer CE durations, the S-HNSS sample showed sustained heavy surface roughening and more obvious boundary protrusions compared to those at the initial CE period (i.e., 1 h, Fig. 5(c)), as shown in Fig. 6(b). The high-magnification image of the eroded surface (Fig. 6(d)) showed that some micro-cracks initiated along the slip lines inside the grains and protrusions formed at the grain boundaries. The mass loss started at the slip lines inside the grains, especially at the intersections of these slip lines. After 5 h of the CE test, severe mass damage was observed in the case of the wrought HNSS sample, as shown in Fig. 7(a). This is consistent with the mass loss rate shown in Fig. 4(b). At CE durations longer than 4 h, the mass loss rate of the wrought HNSS accelerated. Not only its entire original surface was destroyed, large and deep craters and cracks were also formed by collapsing of the cavitation bubbles, and their coalescence led to the tearing-off of the massive materials. Compared to that of the wrought HNSS sample, the S-HNSS sample suffered less mass loss with some original surface preserved, as shown in Fig. 7(b). Fig. 8 shows the morphologies of the eroded surface of the wrought HNSS and S-HNSS samples subjected to CE for 8 h. Although both the wrought HNSS and S-HNSS samples showed similar features after 5 h of the CE test, the size of the craters in the S-HNSS sample was relatively smaller.
Fig. 2. XRD analyses of the wrought HNSS and S-HNSS.
Fig. 3. Load-displacement plots of the wrought HNSS and S-HNSS samples.
Table 2 Indentation parameters derived from the load-displacement curves in Fig. 3.
HNSS S-HNSS
hmax (nm)
hr (nm)
ηh
Hd (GPa)
422.1 377.2
358.3 268.3
0.15 0.29
208.9 120.1
indentation parameters are given in Table 2. The higher ηh value of the S-HNSS suggests that it has better elastic properties than the wrought HNSS sample. The hardness of the wrought HNSS and S-HNSS samples, as calculated using Eqs. (2) and (3), was 208.9 and 120.1 GPa, respectively. This indicates that the wrought HNSS sample has a higher resistance to plastic deformation than the S-HNSS sample. Fig. 4(a) shows the cumulative mass loss as a function of time for the wrought HNSS and S-HNSS under CE conditions. The mass loss of the wrought HNSS sample was larger than that of the S-HNSS sample. After CE for 8 h, the total cumulative mass loss of the wrought HNSS sample was 5.79 mg cm-2, which is twice of that for the S-HNSS sample (2.92 mg cm-2). The evolution of the cumulative mass loss rates of the wrought HNSS and S-HNSS samples is shown in Fig. 4(b). It can be clearly observed that their mass loss rates increased with the CE time, 72
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 4. Cumulative mass loss (a) and mass loss rate (b) as a function of time for the wrought HNSS and S-HNSS under cavitation conditions.
Fig. 5. SEM micrographs of the eroded surface of the wrought HNSS (a), (c) and the S-HNSS (b), (d) after CE of 1 h.
3.4. Cross-sectional observation of the eroded samples
3.5. Micro-hardness measurements in cross section
The cross-sections of the wrought HNSS and S-HNSS samples after the CE test for 8 h are shown in Fig. 9. In the case of the wrought HNSS sample, the micro-cracks (or/and voids) were initiating at the white Cr2N secondary phase precipitated at the grain boundary. The appearance of the larger cracks or voids indicates the subsequent material loss, as shown in Fig. 9(a). Similar corrosion phenomenon was observed in Fe-18Cr-12Mn-XN alloys reported by P. Behjati [2]. No significant difference was observed in the cross-sectional images of the wrought HNSS and S-HNSS, as shown in Fig. 9(b).
To investigate the effect of solution treatment on the work hardening behaviors of the wrought HNSS and S-HNSS, micro-hardness measurements were performed on the cross-sections of these samples at ~20 µm underneath the CE surface at different CE times. Both the wrought HNSS and S-HNSS samples experienced work hardening and showed similar work-hardening rates. The micro-hardness of the samples increased with an increase in the CE duration, indicating that the work hardening behavior of the wrought HNSS and S-HNSS samples was more pronounced at longer CE times, as shown in Fig. 10. 73
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 6. SEM micrographs of the eroded surface of the wrought HNSS (a), (c) and the S-HNSS after (b), (d) CE for 3 h.
4. Discussion
sample had a larger total grain boundary area to prevent the movement of dislocations. An increase in the dislocation density in grains induces work hardening. This is consistent with the results reported by Yuji Ikegami [43], who investigated the effect of finish-forging temperature on the strength and toughness of an Fe-20Cr-15Mn-4Ni-2Mo-0.655 N HNSS and found that the HNSS consisted of slightly elongated fine austenite grains with high dislocation densities at the finish-temperature of 800 °C. When the HNSS was heated to 1150 °C, the dissolution of the Cr2N precipitates and recrystallization of grains occurred, leading to a decrease in the pinning force of the dislocations with the subsequent coarsening of grains. Therefore, the main softening mechanism during the solution treatment is expected to reduce the dislocation density via the recovery process and grain coarsening.
4.1. Effect of solution treatment on mechanical properties The results of the nanoindentation test showed that the hardness of the wrought HNSS sample decreased but with a substantial increase in the elasticity by the solution treatment. The decrease in the hardness of the HNSS by the solution treatment is consistent with the results reported by Ikegami [43]. From the microstructures shown in Fig. 1 and the nanoindentation test results shown in Fig. 3, it can be inferred that the higher hardness value of the wrought HNSS can be attributed to the combination of Hall-Petch hardening due to the grain size [44–46] and the work hardening due to the hot rolling [47]. The wrought HNSS
Fig. 7. SEM micrographs of the eroded surface of (a) the wrought HNSS and (b) the S-HNSS after CE for 5 h. 74
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 8. SEM micrographs of the eroded surfaces of (a), (c) and (e) the wrought HNSS and (b), (d) and (f) the S-HNSS after CE for 8 h.
material with a higher hardness is superior to that of the material with a lower hardness. It is obvious that for the materials with super elasticity and high hardening rates such as NiTi [55–57] and HNSSs, the hardness cannot be the only parameter used to evaluate the CE resistance. The solution treatment of the HNSS increased its CE resistance significantly (to about double). The CE resistance of the wrought HNSS was lower than that of the S-HNSS, though the hardness of the former was higher. This is consistent with the changes in the elastic properties and a higher elastic recovery ratio in the nanoindentation tests, Fig. 3. The increase in the CE resistance of the HNSS after the solution treatment can be attributed to the increase in the elastic energy returned to the environment and a decrease in the amount of the plastic energy absorbed by the sample. As a result, the specimen was plastically loaded to a lesser extent and showed a better resistance to plastic deformation [41]. Planar glide is the dominant plastic deformation mode in incubation period and is related to the SFE of a material [58,59]. It has been
4.2. Structural changes during CE On the basis of the images of the CE morphologies of the samples, it can be stated that the CE damage in this study was a deformationcontrolled process. The CE resistance of a material in general is related to its ability to absorb energy [48,49]. Several mechanisms of energy absorption in the HNSSs have been proposed. These include cold working because of an increase in the dislocation density, planar dislocation gliding, twinning [17], or phase transformations in metastable austenitic alloys [50–52]. The austenite in the HNSSs is quite stable [53] and the strain-induced γ (austenite) → α´ (strain - induced martensite with a b.c.c. crystal structure) or γ → ε (strain - induced martensite with an h.c.p. crystal structure) transformations can be restrained under CE conditions [14]. Therefore, there must be other mechanisms to consume or absorb the energy produced by the collapse of cavitation bubbles. Liu [54] reported that the CE resistance of 75
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 9. Cross-sectional images of (a) the wrought HNSS and (b) the S-HNSS after 8 h of erosion.
dislocations and induced plastic deformation at these sites leading to the undulation of the grain boundaries, as shown in Fig. 5. As the CE process proceeded, the plastic deformation became intense, more twins and grain boundaries were extruded, leading to a clearer appearance of the grain boundaries, as shown in Fig. 6(a) and (b). At the same time, as the number of the slip lines increased and the formation of CE damages initiated at these sites, twin boundaries also began to form. These CE damages on the twin boundaries and slip lines then grew and coalesced. Since the slip lines were close, the CE damages formed on the neighboring slip lines could easily merge, and they appeared to grow across these linear deformation features. With an increase in the CE time, the coalescence of these damages led to the drop of small chunks of the material and the appearance of a grainy damage morphology, as shown in Figs. 7 and 8. 5. Conclusions Fig. 10. Micro-hardness of the tested samples as a function of the CE time at 20 µm underneath the surface.
The effect of solution treatment on the CE behavior of an HNSS was investigated. The solution treatment showed a significant effect on the CE behavior of the HNSS. After the solution treatment at 1150 °C and water quenching, the grain size of the HNSS increased and the precipitated phase disappeared, leading to a decrease in the hardness from 299.5 to 285.1 Hv and a decrease of up to 50% in the cumulative mass loss during the CE test. The increase in the work hardening of the surface layer and the improvement of the elastic properties was beneficial to resist the energy induced during the CE, leading to an increase in the CE resistance.
reported that the addition of nitrogen to austenitic steels decreases their SFE [60–62]. Therefore, the tendency of strain hardening of the highnitrogen austenitic steels is enormously high [12]. The change in the hardness of the material due to the action of load induced by CE reflects the macroscopic changes in its structure [14]. Especially, mechanical twins are easy to form under the repeated attack-induced CE. The impact energy induced by CE can be used for the formation of mechanical twins. In addition, the appearance of mechanical twins tantamount to refined austenitic grains was observed. Thus, the CE resistance was enhanced. Generally, the CE properties of the material depend on its elasticity and work hardening rate. After the CE test for 8 h, the hardness of the wrought HNSS increased from 299.5 to 481.3 Hv, while that of the SHNSS increased from 285.1 to 475.2 Hv. The work hardening rate of the wrought HNSS was comparable to that of the S-HNSS. However, the CE resistance of the S-HNSS was 1.5 times higher than that of the wrought HNSS. Owing to its improved elastic properties which favors erosion energy consumption, the S-HNSS showed higher erosion resistance than the wrought HNSS.
Acknowledgements We express our gratitude to the financial support of the National Natural Science Foundation of China (Nos.51401092, 51409129 and 51304041), Science and Technology Support Program of Jiangsu Province (BE2017143), Project of Science and Technology Development (No. 2017GDASCX-0117), Science and Technology Planning Project of Guangdong Province (No. 2017A070701029) and Dr. S. T. Dong for the nanoindentation tests. References [1] S. Hertzman, The influence of nitrogen on microstructure and properties of highly alloyed stainless steel welds, ISIJ Int. 41 (2001) 580–589. [2] P. Behjati, A. Kermanpur, A. Najafizadeh, Influence of nitrogen alloying on properties of Fe318Cr312Mn3XN austenitic stainless steels, Mater. Sci. Eng. A 588 (2013) 43–48. [3] F.Y. Dong, P. Zhang, J.C. Pang, Y.B. Ren, K. Yang, Z.F. Zhang, Strength, damage and fracture behaviors of high-nitrogen austenitic stainless steel processed by highpressure torsion, Scr. Mater. 96 (2015) 5–8. [4] H.B. Li, Z.H. Jiang, H. Feng, S.C. Zhang, L. Li, P.D. Han, R.D.K. Misra, J.Z. Li,
4.3. The CE damage mechanism of the HNSS On the basis of the evolution of the CE corrosion process (Figs. 5–8), the CE damage of the HNSS can be summarized as follows: at the beginning of the CE action, plastic deformation in the form of slip lines and deformation twins occurred on the surface. During the deformation process, the grain boundaries acted as barriers to prevent the motion of 76
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
[5]
[6] [7]
[8]
[9]
[10] [11]
[12]
[13]
[14] [15]
[16]
[17]
[18] [19]
[20]
[21] [22]
[23]
[24]
[25]
[26]
[27]
[28] [29]
[30]
[31] [32]
[33] K.H. Lo, C.T. Kwok, K.Y. Wang, W.J. Ai, Implications of solution treatment on cavitation erosion and corrosion resistances and synergism of austenitic stainless steel, Wear 392–393 (2017) 159–166. [34] J. Marteau, S. Bouvier, M. Bigerelle, Review on numerical modeling of instrumented indentation tests for elastoplastic material behavior identification, Arch Comput. Method Eng. 22 (2015) 577–593. [35] K. Kowalczyk-Gajewska, K. Sztwiertnia, J. Kawałko, K. Wierzbanowski, M. Wronski, K. Frydrych, S. Stupkiewicz, H. Petryk, Texture evolution in titanium on complex deformation paths: experiment and modelling, Mater. Sci. Eng. A 637 (2015) 251–263. [36] Y.X. Qiao, X. Cai, J. Cui, Y.B. Tang, H.B. Li, Z.H. Jiang, Cavitation erosion properties of a nickel-free high-nitrogen Fe-Cr-Mn-N stainless steel, Mater. Tehnol. 51 (2017) 933–938. [37] B.D. Cullity, Elements of X–Ray Diffraction, 2nd ed, Addison–Wesley Publishing, Reading, MA, 1978. [38] K.H. Kim, G. Chahine, J.P. Franc, et al., Advanced Experimental and Numerical Techniques for Cavitation Erosion Prediction, Springer, Netherlands, 2014. [39] K.J.L. Martin, A. Karimi, Cavitation erosion of materials, Int. Met. Rev. 31 (1986) 1–26. [40] M. Dular, A. Osterman, Pit clustering in cavitation erosion, Wear 265 (2008) 811–820. [41] C.M. Garzón, J.F. Dos Santos, H. Thomas, A.P. Tschiptschin, Cavitation erosion resistance of a high temperature gas nitrided duplex stainless steel in substitute ocean water, Wear 259 (2015) 145–153. [42] A.K. Krella, A. Krupa, Effect of cavitation intensity on degradation of X6CrNiTi1810 stainless steel, Wear 408–409 (2018) 180–189. [43] Y. Ikegami, R. Nemot, Effect of Thermo-mechanical of high-nitrogen containing treatment Cr-Mn-Ni on mechanical properties austenitic stainless steels, ISIJ Int. 36 (1996) 855–861. [44] L. Liu, Y.F. Jia, F.Z. Xuan, Gradient effect in the waved interfacial layer of 304L/ 533B bimetallic plates induced by explosive welding, Mater. Sci. Eng. A 704 (2017) 493–502. [45] N.J. Petch, The cleavage strength of polycrystals, J. Iron Steel Inst. 174 (1953) 25–28. [46] H. Somekawa, T. Mukai, Hall-petch relation for deformation twinning in solid solution magnesium alloys, Mater. Sci. Eng. A 561 (2013) 378–385. [47] L.M. Kaputkina, V.G. Prokoshkina, Martensitic transformations and martensite structure in thermomechanically strengthened high-nitrogen steels, Mater. Sci. Eng. A 438–440 (2006) 228–232. [48] Z.B. Zheng, Y.G. Zheng, W.H. Sun, J.Q. Wang, Effect of heat treatment on the structure, cavitation erosion and erosion-corrosion behavior of Fe-based amorphous coatings, Tribol. Int. 90 (2015) 393–403. [49] S. Lavigne, F. Pougoum, S. Savoie, L. Martinu, J.E. Klemberg-Sapieha, R. Schulz, Cavitation erosion behavior of HVOF caviTec coatings, Wear 386–387 (2017) 90–98. [50] W.T. Fu, Microstructure changes during cavitation erosion for a steel with metastable austenite, J. Mater. Sci. Technol. 16 (2000) 546–548. [51] Z. Xiaojun, L. Procopiak, N. Souza, A. d'Oliveira, Phase transformation during cavitation erosion of a Co stainless steel, Mater. Sci. Eng. A 358 (2003) 199–204. [52] M.C. Park, K.N. Kim, G.S. Shin, S.J. Kim, Effects of strain induced martensitic transformation on the cavitation erosion resistance and incubation time of Fe-Cr-NiC Alloys, Wear 274–275 (2012) 28–33. [53] Y.X. Qiao, Y.G. Zheng, X.Q. Wu, W. Ke, K. Yang, Z.H. Jiang, Cavitation erosion resistance of high nitrogen stainless steel in comparison with low N content CrMnN stainless steel, Tribol. Mater. Surf. Inter. 1 (2013) 165–172. [54] W. Liu, Y.G. Zheng, Z.M. Yao, X.Q. Wu, W. Ke, Cavitation erosion of 20SiMn and 0Cr13Ni5Mo steels in distilled water with and without sand, Acta Metall. Sin. 37 (2001) 197–201. [55] K.Y. Chiu, F.T. Cheng, H.C. Man, Hydrogen effect on the cavitation erosion resistance of AISI 316L stainless steel laser surface-modified with NiTi, Mater. Lett. 61 (2007) 239–243. [56] K.Y. Chiu, F.T. Cheng, H.C. Man, Cavitation erosion resistance of AISI 316L stainless steel laser surface-modified with NiTi, Mater. Sci. Eng. A 392 (2005) 348–358. [57] W. Liu, Y.G. Zheng, C.S. Liu, Z.M. Yao, Cavitation erosion characteristics of a NiTi alloy, Metall. Mater. Trans. A 35 (2004) 356–362. [58] J.B. Vogt, T. Magnin, J. Foct, Effective stresses and microstructure in cyclically deformed 316L austenitic stainless steel: effect of temperature and nitrogen content, Fatigue Fract. Eng. M 16 (2010) 555–564. [59] J. Polák, F. Fardoun, S. Degallaix, Analysis of the hysteresis loop in stainless steels II. austenitic-ferritic duplex steel and the effect of nitrogen, Mater. Sci. Eng. A 297 (2001) 154–161. [60] R. Fawley, M.A. Quader, R.A. Dodd, Compositional effects on the deformation modes, annealing twin frequencies and stacking fault energies of austenitic stainless steels, Trans. Metall. Soc. AIME 242 (1968) 771–776. [61] P.R. Swann, Dislocation substructure vs transgranular stress corrosion susceptibility of single phase alloys, Corrosion 19 (1963) 102–114. [62] R.E. Schramm, R.P. Reed, Stacking fault energies of seven commercial austenitic stainless steels, Metall. Trans. A 6 (1975) 1345–1351.
Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel, Mater. Des. 84 (2015) 291–299. M. Moallemi, A. Zarei-Hanzaki, H.S. Baghbadorani, Evolution of microstructure and mechanical properties in a cold deformed nitrogen bearing TRIP-assisted duplex stainless steel after reversion annealing, Mater. Sci. Eng. A 683 (2017) 83–89. A. Di Schino, M.G. Mecozzi, M. Barteri, Solidification mode and residual ferrite in low-Ni austenitic stainless steels, J. Mater. Sci. 35 (2000) 375–380. H.B. Li, Z.H. Jiang, Y. Cao, Z.R. Zhang, Fabrication of high nitrogen austenitic stainless steels with excellent mechanical and pitting corrosion properties, Int. J. Miner. Metall. Mater. 16 (2009) 387–392. M. Talha, C.K. Behera, O.P. Sinha, Effect of nitrogen and cold working on structural and mechanical behavior of Ni-free nitrogen containing austenitic stainless steels for biomedical applications, Mater. Sci. Eng. C 47 (2015) 196–203. L.M. Kaputkina, V.G. Prokoshkina, Martensitic transformations and martensite structure in thermomechanically strengthened high-nitrogen steels, Mater. Sci. Eng. A 438–440 (2006) 228–232. Y.X. Qiao, Y.G. Zheng, W. Ke, P.C. Okafor, Electrochemical behaviour of high nitrogen stainless steel in acidic solutions, Corros. Sci. 51 (2009) 979–986. Y.X. Qiao, X. Cai, J. Cui, H.B. Li, Passivity and semiconducting behavior of a high nitrogen stainless steel in acidic NaCl solution, Adv. Mater. Sci. Eng. (2016) (No.6065481). P. Wan, Y.B. Ren, B.C. Zhang, K. Yang, Effect of nitrogen on biocorrosion behavior of high nitrogen Nickel-free stainless steel in different simulated body fluids, Mater. Sci. Eng. C 32 (2012) 510–516. Y.X. Qiao, Z.H. Tian, X. Cai, J. Chen, Y.X. Wang, Q.N. Song, H.B. Li, Cavitation erosion behaviors of a nickel-free high-nitrogen stainless steel, Tribol. Lett. 67 (2019) 1–9. W.T. Fu, Y.Z. Zheng, X.K. He, Resistance of a high nitrogen austenitic steel to cavitation erosion, Wear 249 (2001) 788–791. W. Liu, Y.G. Zheng, C.S. Liu, Z.M. Yao, W. Ke, Cavitation erosion behavior of Cr-MnN stainless steels in comparison with 0Cr13Ni5Mo stainless steel, Wear 254 (2003) 713–722. G. Bregliozzia, A. Di Schinob, S.I.-U. Ahmeda, J.M. Kenny, H. Haefkea, Cavitation wear behaviour of austenitic stainless steels with different grain sizes, Wear 258 (2005) 503–510. D.H. Mesa, C.M. Garzón, A.P. Tschiptschin, Influence of cold-work on the cavitation erosion resistance and on the damage mechanisms in high-nitrogen austenitic stainless steels, Wear 271 (2011) 1372–1377. P. Niederhofer, S. Huth, Cavitation erosion resistance of high interstitial CrMnCN austenitic stainless steels, Wear 301 (2013) 457–466. Y.X. Qiao, S. Wang, B. Liu, Y.G. Zheng, H.B. Li, Z.H. Jiang, Synergistic effect of corrosion and cavitation erosion of high nitrogen stainless steel, Acta Metall. Sin. 52 (2016) 233–240. J. Li, Y.X. Yang, Y.B. Ren, J.H. Dong, K. Yang, Effect of cold deformation on corrosion fatigue behavior of nickel-free high nitrogen austenitic stainless steel for coronary stent application, J. Mater. Sci. Technol. 34 (2018) 660–665. J.W. Simmons, Influence of nitride (Cr2N) precipitation on the plastic flow behavior of high-nitrogen austenitic stainless steel, Scr. Metall. 32 (1995) 265–270. H. Lee, S.J. Kim, Y.C. Jung, Crystallographic details of precipitates in Fe-22Cr-21Ni6Mo-(N) superaustenitic stainless steels aged at 900 °C, Metall. Mater. Trans. A 31 (2000) 1713–1723. R.D. Knutsen, C.L. Lang, J.A. Basson, Discontinuous cellular precipitation in a CrMn-N steel with niobium and vanadium additions, Acta Mater. 52 (2004) 2407–2417. H.B. Li, Z.H. Jiang, H. Feng, Q.F. M, D.P. Zhan, Aging precipitation behavior of 18Cr-16Mn-2Mo-1.1N high nitrogen austenitic stainless steel and its influences on mechanical properties, J. Iron Steel Res. Int. 19 (2012) 43–51. N. Sathirachinda, R. Pettersson, J. Pan, Depletion effects at phase boundaries in 2205 duplex stainless steel characterized with SKPFM and TEM/EDS, Corros. Sci. 51 (2009) 1850–1860. Y.H. Yang, B. Yan, J. Wang, J.L. Yin, The influence of solution treatment temperature on microstructure and corrosion behavior of high temperature ageing in 25% Cr duplex stainless steel, J. Alloy Compd. 509 (2011) 8870–8879. M.Q. Ou, Y.C. Ma, X.C. Hao, B.F. Wan, T. Liang, K. Liu, Effect of solution annealing on microstructure and mechanical properties of a Ni-Cr-W-Fe alloy, J. Mater. Sci. Technol. 33 (2017) 1300–1307. C.T. Kwok, H.C. Man, F.T. Cheng, Cavitation erosion and damage mechanisms of alloys with duplex structures, Mater. Sci. Eng. A 242 (1998) 108–120. Z.Y. Wang, J.H. Zhu, Correlation of martensitic transformation and surface mechanical behavior with cavitation erosion resistance for some iron-based alloys, Wear 256 (2004) 1208–1213. R.B. Nair, H.S. Arora, S. Mukherjee, S. Singh, H. Singh, H.S. Grewal, Exceptionally high cavitation erosion and corrosion resistance of a high entropy alloy, Ultrason. Sonochem. 41 (2018) 252–260. F.T. Cheng, P. Shi, H.C. Man, Correlation of cavitation erosion resistance with indentation-derived properties for a NiYi alloy, Scr. Mater. 45 (2001) 1083–1089. I. Mitelea, M.L. Micu, I. Bordeasu, C.M. Crãciunescu, Cavitation erosion of sensitized UNS S31803 duplex stainless steels, J. Mater. Eng. Perform. 25 (2016) 1939–1944.
77
Contents lists available at ScienceDirect
Wear journal homepage: www.elsevier.com/locate/wear
Effect of solution treatment on cavitation erosion behavior of high-nitrogen austenitic stainless steel ⁎
T
⁎
Yanxin Qiaoa, , Jian Chena, , Huiling Zhoua, Yuxin Wanga, Qining Songb, Huabing Lic, Zhibin Zhengd a
School of Materials Science and Engineering, Jiangsu University of Science and Technology, Zhenjiang 212003, China College of Mechanical and Electrical Engineering, Hohai University, Changzhou 213022, China c School of Metallurgy, Northeastern University, Shenyang 11004, China d Guangdong Key Laboratory for Technology and Application of Metal Toughening, Guangdong Institute of Materials and Processing, Guangzhou 510650, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: High-nitrogen austenitic stainless steel Solution treatment Cavitation erosion Elastic property
The effect of solution treatment on the cavitation erosion behavior of a Fe-18Cr-16Mn-2Mo-0.66 N high-nitrogen austenitic stainless steel (HNSS) exposed in distilled water up to 8 h was investigated using magnetostrictive cavitation, micro-hardness, mass loss measurements, scanning electron microscope (SEM) and X-ray diffraction (XRD). The results showed that the solution treatment slightly decreased the hardness of the HNSS but significantly increased its elasticity. The HNSS with an improved elastic property was shown to have a higher resistance to the absorption of energy produced by cavitation erosion, and hence a higher cavitation erosion resistance. The cavitation erosion damage of the HNSS in distilled water was found to be a deformation-controlled process which limited by the formation and development of plastic deformation during cavitation erosion tests.
1. Introduction Austenitic stainless steels have been widely used in various industrial fields because of their excellent mechanical and corrosion properties. However, the low yield strength limits their widespread application. In order to overcome this limitation, nitrogen is used as an additive in austenitic stainless steels [1–5]. High-nitrogen nickel-free stainless steels (HNSSs) have developed and gained immense attention for various industrial applications owing to their low cost [6], high strength and ductility [7], high work-hardening ability [8], high stressinduced martensite transformation ability [9], high corrosion resistance [10,11], and good biocompatibility [12]. Moreover, these alloys also show excellent cavitation erosion (CE) resistance [13–19], which depends on their microstructure and mechanical properties. Nitrogencontaining austenitic stainless steels also show high yield strength, ductility, and fracture toughness [1,3,20]. However, when exposed to temperatures higher than 900 °C for a short duration (i.e., 2 h), their mechanical properties would deteriorate due to the precipitation of Cr2N secondary phase and the formation of (Mn, Cr)-oxides in the HNSSs and nitrogen-containing austenitic stainless steels [21–24]. Solution treatment is widely used to homogenize metallic alloys, which can reduce the precipitation of the secondary phases, optimize ⁎
the microstructure, and improve the strength [25–27]. Although solution treatment significantly affects the microstructure, grain size, and mechanical properties of the alloys, the effect of solution treatment on the CE property of the metallic alloys still needs further investigations [28–30]. The solution treatment of UNS S30400 stainless steel decreased its CE resistance because of the increased grain size and ductility [28]. However, the effect of solution treatment on the CE behavior of S32760 duplex stainless steel was found be to minimal since the mechanical properties and grain size of the S32760 were not significantly affected by solution treatment [28]. Cheng et al. [31] investigated the CE behavior of the solution-treated and aged (aged after solution treatment) NiTi (50.8 at% Ni) and found that the CE resistance of the solution-treated sample was higher than that of the samples aged at temperatures in the range from 200 to 600 °C. The sensitization treatment of UNS S31803 duplex stainless steel at 475 °C deteriorated its CE resistance (~ 22% lower than that obtained after the solution annealing treatment) [32]. High-temperature (1300 °C) solution treatment of CrMnNi austenitic stainless steels led to a significant grain growth, which compromised their strength and facilitated the formation of martensite, and thereby deteriorated their CE properties [33]. Therefore, further studies on the CE process and mechanism are required to investigate the effect of heat treatment on the CE behavior of
Corresponding authors. E-mail addresses: [email protected] (Y. Qiao), [email protected] (J. Chen).
https://doi.org/10.1016/j.wear.2019.01.098 Received 2 November 2018; Received in revised form 29 January 2019; Accepted 29 January 2019 Available online 30 January 2019 0043-1648/ © 2019 Elsevier B.V. All rights reserved.
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
where hc is the contact depth at the peak load, P (h) is the corresponding load, ε is a correction index (0.75 for Berkovich indenter), and S is the slope of the unloading curve. The CE test was carried out in distilled water at 20 ± 1 °C. The dimensions of the specimen for the cavitation test and the details of the test were described elsewhere [13]. The CE test was repeated three times for each sample. The microstructures and eroded morphologies of the HNSS samples were observed using scanning electron microscopy (XL30-FEG ESEM). The crystal structures of the treated (solution treatment) and untreated HNSS samples were examined using a D8 Xray diffractionmeter (XRD) with a Cu Ka radiation at 40 kV and 35 mA. To investigate the effect of cavitation on work hardening of the samples, their micro-hardness was measured at a fixed interval using an MH-5 tester with a load of 100 g and a loading time of 15 s. These measurements were performed on the cross-sections of the specimens after CE for 1, 3, 5 and 8 h at a distance of 20 µm below the CE surface for five times. The average value of five measurements was used to represent the hardness of the eroded specimens.
Table 1 Mechanical parameters of the HNSS and S-HNSS.
HNSS S-HNSS
Rm (MPa)
Rp0.2 (MPa)
Hv100
920.4 ± 22.1 855.8 ± 16.3
570.6 ± 15.4 505.2 ± 14.3
299.5 ± 5.3 285.1 ± 4.5
metallic alloys such as the HNSSs. In this study, CE tests were carried out on a nickel-free HNSS (Fe18Cr-16Mn-2Mo-0.66 N) in distilled water using an ultrasonic vibration system. The effect of solution treatment on the CE properties and damage mechanism of this nickel-free HNSS were investigated. The erosion mechanisms were discussed. 2. Experimental details The chemical composition (wt%) of the Fe-18Cr-16Mn-2Mo-0.66 N HNSS used in this study was C 0.048%, Si 0.24%, Mn 15.96%, P 0.004%, S 0.017%, Cr 18.44%, Mo 2.23%, N 0.66%, and Fe balanced. The hot-forged ingots were cut into round bars with a diameter of 20 mm and were labeled as wrought HNSS. Some of the wrought HNSS samples were heated to 1150 °C for 1 h and were then quenched in water. These samples were named as S-HNSS. Tensile tests were carried out on an Instron-type testing machine with a tensile rate of 3 mm s-1 at 25 ± 1 °C. The mechanical properties of the HNSS and S-HNSS, including the surface hardness (Hv100), the tensile Rm and yield Rp0.2 strength are shown in Table 1. To characterize the elastic properties of the HNSS and S-HNSS, nanoindentation curves were acquired using a CSM NHT2 nanoindenter (Anton Paar) at a maximum load of 20 mN with a loading/unloading rate of 40 mN·min-1. In nanoindentation tests, the elastic behavior of a material can be analyzed by its depth recovery ratio (ηh) obtained from the load-displacement curve [34]. The value of ηh is calculated as follows:
ηh =
h max − hr h max
3. Results 3.1. Material characterization Fig. 1 shows the microstructure of the wrought HNSS and S-HNSS. The wrought HNSS was composed of an austenitic phase matrix (marked as 1 in Fig. 1(a)) with twins and homogeneously-distributed fine nitrides Cr2N particles (marked as 2 in Fig. 1(a)) formed during the solidification process [36]. The S-HNSS exhibited a single austenite phase (marked as 3 in Fig. 1(b)) with twins distributing inside the austenite grains. The grain size of the S-HNSS sample was larger (~25 µm) than that of the wrought HNSS sample (~8 µm). Fig. 2 shows the XRD patterns of the wrought HNSS and S-HNSS samples. Although both the wrought HNSS and S-HNSS were composed of a single γ phase with a face centered cubic (f.c.c.) structure, the characteristic peaks of the γ phase in the wrought HNSS were broader than those of the S-HNSS. According to the Scherrer equation [37], the grain size of a material is inversely proportional to the width of its peaks. Hence, the grain size of the wrought HNSS was smaller than that of the S-HNSS. This is consistent with the microscopic results shown in Fig. 1.
(1)
where hmax is the maximum penetration depth and hr is the residual depth after unloading. The hardness of stainless steels is defined as the ratio of the peak indentation load (Fmax) to the hardness impression projected area (Ac) and is calculated using Eqs. (2) and (3) as follows [35]:
F Fmax Hd = max = Ac 26.43hc2
hc = h max − ε
P (h) S
3.2. Nanoindentation test Fig. 3 shows the nanoindentation load-displacement curves (P-h) of the wrought HNSS and S-HNSS samples. The maximum indentation depth (hmax) at a load of 20 mN of the wrought HNSS sample (422.1 nm) was larger than that of the S-HNSS (377.2 nm). The other
(2) (3)
Fig. 1. The microstructure of the wrought HNSS (a) and S-HNSS (b). 71
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
and the mass loss rate of the wrought HNSS sample was larger than that of the S-HNSS sample. Both the wrought HNSS and S-HNSS samples demonstrated a plateau period and acceleration period [38]. Although the plateau periods for both the samples were observed during the initial 4 h, the mass loss rate of the wrought HNSS in this period was 3.4 times of that of the S-HNSS, indicating that the S-HNSS showed a high cavitation resistance in distilled water. Once the acceleration period of CE started, the erosion damage spread over the entire area of the specimen [39], leading to a rapid increase in the mass-loss rate. During the accumulation period, the slopes of the weight loss curves of the wrought HNSS and S-HNSS were 2.0 × 10-3 and 1.6 × 10-3 mg cm-2 h-1, respectively, indicating that the CE resistance of the HNSS was improved by the solution treatment.
3.3. Morphologies of the eroded surfaces Fig. 5 shows the SEM micrographs of the wrought HNSS and S-HNSS samples after erosion for 1 h. As shown in Fig. 5(a) and (c), the wrought HNSS sample experienced deformation in the presence of slip lines and deformed twins inside the grains, and surface undulations. In addition, the wrought HNSS suffered from material loss at the grain and twin boundaries. Some small holes with either circular or angular shapes were also observed on the eroded surface (black arrows in Fig. 5(c)). Taking into account the geometry of these holes, it is most likely that they were formed by the spalling of the Cr2N precipitation during the CE test. This suggests that the damage induced by CE was preferentially initiated at the grain boundaries, slip lines, and precipitates [40]. Although the S-HNSS sample exhibited plastic deformation, its material loss was minor, as shown in Fig. 5(b) and (d). This is consistent with the mass loss results shown in Fig. 4. This further demonstrated that solution treatment can improve the CE resistance of the HNSSs. The damaged surfaces of the wrought HNSS and S-HNSS samples after erosion for 3 h are shown in Fig. 6. With an increase in the CE duration, both the specimens experienced more pronounced deformation. The wrought HNSS exhibited obvious cavitation damage at the grain and twin boundaries, as shown in Fig. 6(a) and (c), which started to spread into the interior of the grains. Although some grains underwent intense plastic deformation with the establishment of multiple slips and initiation of micro-cracks inside the grains, some grains still remained intact and smooth as the original polished surface. This can be attributed to the different values of the resolved shear stress inside the grains [41,42]. At longer CE durations, the S-HNSS sample showed sustained heavy surface roughening and more obvious boundary protrusions compared to those at the initial CE period (i.e., 1 h, Fig. 5(c)), as shown in Fig. 6(b). The high-magnification image of the eroded surface (Fig. 6(d)) showed that some micro-cracks initiated along the slip lines inside the grains and protrusions formed at the grain boundaries. The mass loss started at the slip lines inside the grains, especially at the intersections of these slip lines. After 5 h of the CE test, severe mass damage was observed in the case of the wrought HNSS sample, as shown in Fig. 7(a). This is consistent with the mass loss rate shown in Fig. 4(b). At CE durations longer than 4 h, the mass loss rate of the wrought HNSS accelerated. Not only its entire original surface was destroyed, large and deep craters and cracks were also formed by collapsing of the cavitation bubbles, and their coalescence led to the tearing-off of the massive materials. Compared to that of the wrought HNSS sample, the S-HNSS sample suffered less mass loss with some original surface preserved, as shown in Fig. 7(b). Fig. 8 shows the morphologies of the eroded surface of the wrought HNSS and S-HNSS samples subjected to CE for 8 h. Although both the wrought HNSS and S-HNSS samples showed similar features after 5 h of the CE test, the size of the craters in the S-HNSS sample was relatively smaller.
Fig. 2. XRD analyses of the wrought HNSS and S-HNSS.
Fig. 3. Load-displacement plots of the wrought HNSS and S-HNSS samples.
Table 2 Indentation parameters derived from the load-displacement curves in Fig. 3.
HNSS S-HNSS
hmax (nm)
hr (nm)
ηh
Hd (GPa)
422.1 377.2
358.3 268.3
0.15 0.29
208.9 120.1
indentation parameters are given in Table 2. The higher ηh value of the S-HNSS suggests that it has better elastic properties than the wrought HNSS sample. The hardness of the wrought HNSS and S-HNSS samples, as calculated using Eqs. (2) and (3), was 208.9 and 120.1 GPa, respectively. This indicates that the wrought HNSS sample has a higher resistance to plastic deformation than the S-HNSS sample. Fig. 4(a) shows the cumulative mass loss as a function of time for the wrought HNSS and S-HNSS under CE conditions. The mass loss of the wrought HNSS sample was larger than that of the S-HNSS sample. After CE for 8 h, the total cumulative mass loss of the wrought HNSS sample was 5.79 mg cm-2, which is twice of that for the S-HNSS sample (2.92 mg cm-2). The evolution of the cumulative mass loss rates of the wrought HNSS and S-HNSS samples is shown in Fig. 4(b). It can be clearly observed that their mass loss rates increased with the CE time, 72
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 4. Cumulative mass loss (a) and mass loss rate (b) as a function of time for the wrought HNSS and S-HNSS under cavitation conditions.
Fig. 5. SEM micrographs of the eroded surface of the wrought HNSS (a), (c) and the S-HNSS (b), (d) after CE of 1 h.
3.4. Cross-sectional observation of the eroded samples
3.5. Micro-hardness measurements in cross section
The cross-sections of the wrought HNSS and S-HNSS samples after the CE test for 8 h are shown in Fig. 9. In the case of the wrought HNSS sample, the micro-cracks (or/and voids) were initiating at the white Cr2N secondary phase precipitated at the grain boundary. The appearance of the larger cracks or voids indicates the subsequent material loss, as shown in Fig. 9(a). Similar corrosion phenomenon was observed in Fe-18Cr-12Mn-XN alloys reported by P. Behjati [2]. No significant difference was observed in the cross-sectional images of the wrought HNSS and S-HNSS, as shown in Fig. 9(b).
To investigate the effect of solution treatment on the work hardening behaviors of the wrought HNSS and S-HNSS, micro-hardness measurements were performed on the cross-sections of these samples at ~20 µm underneath the CE surface at different CE times. Both the wrought HNSS and S-HNSS samples experienced work hardening and showed similar work-hardening rates. The micro-hardness of the samples increased with an increase in the CE duration, indicating that the work hardening behavior of the wrought HNSS and S-HNSS samples was more pronounced at longer CE times, as shown in Fig. 10. 73
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 6. SEM micrographs of the eroded surface of the wrought HNSS (a), (c) and the S-HNSS after (b), (d) CE for 3 h.
4. Discussion
sample had a larger total grain boundary area to prevent the movement of dislocations. An increase in the dislocation density in grains induces work hardening. This is consistent with the results reported by Yuji Ikegami [43], who investigated the effect of finish-forging temperature on the strength and toughness of an Fe-20Cr-15Mn-4Ni-2Mo-0.655 N HNSS and found that the HNSS consisted of slightly elongated fine austenite grains with high dislocation densities at the finish-temperature of 800 °C. When the HNSS was heated to 1150 °C, the dissolution of the Cr2N precipitates and recrystallization of grains occurred, leading to a decrease in the pinning force of the dislocations with the subsequent coarsening of grains. Therefore, the main softening mechanism during the solution treatment is expected to reduce the dislocation density via the recovery process and grain coarsening.
4.1. Effect of solution treatment on mechanical properties The results of the nanoindentation test showed that the hardness of the wrought HNSS sample decreased but with a substantial increase in the elasticity by the solution treatment. The decrease in the hardness of the HNSS by the solution treatment is consistent with the results reported by Ikegami [43]. From the microstructures shown in Fig. 1 and the nanoindentation test results shown in Fig. 3, it can be inferred that the higher hardness value of the wrought HNSS can be attributed to the combination of Hall-Petch hardening due to the grain size [44–46] and the work hardening due to the hot rolling [47]. The wrought HNSS
Fig. 7. SEM micrographs of the eroded surface of (a) the wrought HNSS and (b) the S-HNSS after CE for 5 h. 74
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 8. SEM micrographs of the eroded surfaces of (a), (c) and (e) the wrought HNSS and (b), (d) and (f) the S-HNSS after CE for 8 h.
material with a higher hardness is superior to that of the material with a lower hardness. It is obvious that for the materials with super elasticity and high hardening rates such as NiTi [55–57] and HNSSs, the hardness cannot be the only parameter used to evaluate the CE resistance. The solution treatment of the HNSS increased its CE resistance significantly (to about double). The CE resistance of the wrought HNSS was lower than that of the S-HNSS, though the hardness of the former was higher. This is consistent with the changes in the elastic properties and a higher elastic recovery ratio in the nanoindentation tests, Fig. 3. The increase in the CE resistance of the HNSS after the solution treatment can be attributed to the increase in the elastic energy returned to the environment and a decrease in the amount of the plastic energy absorbed by the sample. As a result, the specimen was plastically loaded to a lesser extent and showed a better resistance to plastic deformation [41]. Planar glide is the dominant plastic deformation mode in incubation period and is related to the SFE of a material [58,59]. It has been
4.2. Structural changes during CE On the basis of the images of the CE morphologies of the samples, it can be stated that the CE damage in this study was a deformationcontrolled process. The CE resistance of a material in general is related to its ability to absorb energy [48,49]. Several mechanisms of energy absorption in the HNSSs have been proposed. These include cold working because of an increase in the dislocation density, planar dislocation gliding, twinning [17], or phase transformations in metastable austenitic alloys [50–52]. The austenite in the HNSSs is quite stable [53] and the strain-induced γ (austenite) → α´ (strain - induced martensite with a b.c.c. crystal structure) or γ → ε (strain - induced martensite with an h.c.p. crystal structure) transformations can be restrained under CE conditions [14]. Therefore, there must be other mechanisms to consume or absorb the energy produced by the collapse of cavitation bubbles. Liu [54] reported that the CE resistance of 75
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
Fig. 9. Cross-sectional images of (a) the wrought HNSS and (b) the S-HNSS after 8 h of erosion.
dislocations and induced plastic deformation at these sites leading to the undulation of the grain boundaries, as shown in Fig. 5. As the CE process proceeded, the plastic deformation became intense, more twins and grain boundaries were extruded, leading to a clearer appearance of the grain boundaries, as shown in Fig. 6(a) and (b). At the same time, as the number of the slip lines increased and the formation of CE damages initiated at these sites, twin boundaries also began to form. These CE damages on the twin boundaries and slip lines then grew and coalesced. Since the slip lines were close, the CE damages formed on the neighboring slip lines could easily merge, and they appeared to grow across these linear deformation features. With an increase in the CE time, the coalescence of these damages led to the drop of small chunks of the material and the appearance of a grainy damage morphology, as shown in Figs. 7 and 8. 5. Conclusions Fig. 10. Micro-hardness of the tested samples as a function of the CE time at 20 µm underneath the surface.
The effect of solution treatment on the CE behavior of an HNSS was investigated. The solution treatment showed a significant effect on the CE behavior of the HNSS. After the solution treatment at 1150 °C and water quenching, the grain size of the HNSS increased and the precipitated phase disappeared, leading to a decrease in the hardness from 299.5 to 285.1 Hv and a decrease of up to 50% in the cumulative mass loss during the CE test. The increase in the work hardening of the surface layer and the improvement of the elastic properties was beneficial to resist the energy induced during the CE, leading to an increase in the CE resistance.
reported that the addition of nitrogen to austenitic steels decreases their SFE [60–62]. Therefore, the tendency of strain hardening of the highnitrogen austenitic steels is enormously high [12]. The change in the hardness of the material due to the action of load induced by CE reflects the macroscopic changes in its structure [14]. Especially, mechanical twins are easy to form under the repeated attack-induced CE. The impact energy induced by CE can be used for the formation of mechanical twins. In addition, the appearance of mechanical twins tantamount to refined austenitic grains was observed. Thus, the CE resistance was enhanced. Generally, the CE properties of the material depend on its elasticity and work hardening rate. After the CE test for 8 h, the hardness of the wrought HNSS increased from 299.5 to 481.3 Hv, while that of the SHNSS increased from 285.1 to 475.2 Hv. The work hardening rate of the wrought HNSS was comparable to that of the S-HNSS. However, the CE resistance of the S-HNSS was 1.5 times higher than that of the wrought HNSS. Owing to its improved elastic properties which favors erosion energy consumption, the S-HNSS showed higher erosion resistance than the wrought HNSS.
Acknowledgements We express our gratitude to the financial support of the National Natural Science Foundation of China (Nos.51401092, 51409129 and 51304041), Science and Technology Support Program of Jiangsu Province (BE2017143), Project of Science and Technology Development (No. 2017GDASCX-0117), Science and Technology Planning Project of Guangdong Province (No. 2017A070701029) and Dr. S. T. Dong for the nanoindentation tests. References [1] S. Hertzman, The influence of nitrogen on microstructure and properties of highly alloyed stainless steel welds, ISIJ Int. 41 (2001) 580–589. [2] P. Behjati, A. Kermanpur, A. Najafizadeh, Influence of nitrogen alloying on properties of Fe318Cr312Mn3XN austenitic stainless steels, Mater. Sci. Eng. A 588 (2013) 43–48. [3] F.Y. Dong, P. Zhang, J.C. Pang, Y.B. Ren, K. Yang, Z.F. Zhang, Strength, damage and fracture behaviors of high-nitrogen austenitic stainless steel processed by highpressure torsion, Scr. Mater. 96 (2015) 5–8. [4] H.B. Li, Z.H. Jiang, H. Feng, S.C. Zhang, L. Li, P.D. Han, R.D.K. Misra, J.Z. Li,
4.3. The CE damage mechanism of the HNSS On the basis of the evolution of the CE corrosion process (Figs. 5–8), the CE damage of the HNSS can be summarized as follows: at the beginning of the CE action, plastic deformation in the form of slip lines and deformation twins occurred on the surface. During the deformation process, the grain boundaries acted as barriers to prevent the motion of 76
Wear 424–425 (2019) 70–77
Y. Qiao, et al.
[5]
[6] [7]
[8]
[9]
[10] [11]
[12]
[13]
[14] [15]
[16]
[17]
[18] [19]
[20]
[21] [22]
[23]
[24]
[25]
[26]
[27]
[28] [29]
[30]
[31] [32]
[33] K.H. Lo, C.T. Kwok, K.Y. Wang, W.J. Ai, Implications of solution treatment on cavitation erosion and corrosion resistances and synergism of austenitic stainless steel, Wear 392–393 (2017) 159–166. [34] J. Marteau, S. Bouvier, M. Bigerelle, Review on numerical modeling of instrumented indentation tests for elastoplastic material behavior identification, Arch Comput. Method Eng. 22 (2015) 577–593. [35] K. Kowalczyk-Gajewska, K. Sztwiertnia, J. Kawałko, K. Wierzbanowski, M. Wronski, K. Frydrych, S. Stupkiewicz, H. Petryk, Texture evolution in titanium on complex deformation paths: experiment and modelling, Mater. Sci. Eng. A 637 (2015) 251–263. [36] Y.X. Qiao, X. Cai, J. Cui, Y.B. Tang, H.B. Li, Z.H. Jiang, Cavitation erosion properties of a nickel-free high-nitrogen Fe-Cr-Mn-N stainless steel, Mater. Tehnol. 51 (2017) 933–938. [37] B.D. Cullity, Elements of X–Ray Diffraction, 2nd ed, Addison–Wesley Publishing, Reading, MA, 1978. [38] K.H. Kim, G. Chahine, J.P. Franc, et al., Advanced Experimental and Numerical Techniques for Cavitation Erosion Prediction, Springer, Netherlands, 2014. [39] K.J.L. Martin, A. Karimi, Cavitation erosion of materials, Int. Met. Rev. 31 (1986) 1–26. [40] M. Dular, A. Osterman, Pit clustering in cavitation erosion, Wear 265 (2008) 811–820. [41] C.M. Garzón, J.F. Dos Santos, H. Thomas, A.P. Tschiptschin, Cavitation erosion resistance of a high temperature gas nitrided duplex stainless steel in substitute ocean water, Wear 259 (2015) 145–153. [42] A.K. Krella, A. Krupa, Effect of cavitation intensity on degradation of X6CrNiTi1810 stainless steel, Wear 408–409 (2018) 180–189. [43] Y. Ikegami, R. Nemot, Effect of Thermo-mechanical of high-nitrogen containing treatment Cr-Mn-Ni on mechanical properties austenitic stainless steels, ISIJ Int. 36 (1996) 855–861. [44] L. Liu, Y.F. Jia, F.Z. Xuan, Gradient effect in the waved interfacial layer of 304L/ 533B bimetallic plates induced by explosive welding, Mater. Sci. Eng. A 704 (2017) 493–502. [45] N.J. Petch, The cleavage strength of polycrystals, J. Iron Steel Inst. 174 (1953) 25–28. [46] H. Somekawa, T. Mukai, Hall-petch relation for deformation twinning in solid solution magnesium alloys, Mater. Sci. Eng. A 561 (2013) 378–385. [47] L.M. Kaputkina, V.G. Prokoshkina, Martensitic transformations and martensite structure in thermomechanically strengthened high-nitrogen steels, Mater. Sci. Eng. A 438–440 (2006) 228–232. [48] Z.B. Zheng, Y.G. Zheng, W.H. Sun, J.Q. Wang, Effect of heat treatment on the structure, cavitation erosion and erosion-corrosion behavior of Fe-based amorphous coatings, Tribol. Int. 90 (2015) 393–403. [49] S. Lavigne, F. Pougoum, S. Savoie, L. Martinu, J.E. Klemberg-Sapieha, R. Schulz, Cavitation erosion behavior of HVOF caviTec coatings, Wear 386–387 (2017) 90–98. [50] W.T. Fu, Microstructure changes during cavitation erosion for a steel with metastable austenite, J. Mater. Sci. Technol. 16 (2000) 546–548. [51] Z. Xiaojun, L. Procopiak, N. Souza, A. d'Oliveira, Phase transformation during cavitation erosion of a Co stainless steel, Mater. Sci. Eng. A 358 (2003) 199–204. [52] M.C. Park, K.N. Kim, G.S. Shin, S.J. Kim, Effects of strain induced martensitic transformation on the cavitation erosion resistance and incubation time of Fe-Cr-NiC Alloys, Wear 274–275 (2012) 28–33. [53] Y.X. Qiao, Y.G. Zheng, X.Q. Wu, W. Ke, K. Yang, Z.H. Jiang, Cavitation erosion resistance of high nitrogen stainless steel in comparison with low N content CrMnN stainless steel, Tribol. Mater. Surf. Inter. 1 (2013) 165–172. [54] W. Liu, Y.G. Zheng, Z.M. Yao, X.Q. Wu, W. Ke, Cavitation erosion of 20SiMn and 0Cr13Ni5Mo steels in distilled water with and without sand, Acta Metall. Sin. 37 (2001) 197–201. [55] K.Y. Chiu, F.T. Cheng, H.C. Man, Hydrogen effect on the cavitation erosion resistance of AISI 316L stainless steel laser surface-modified with NiTi, Mater. Lett. 61 (2007) 239–243. [56] K.Y. Chiu, F.T. Cheng, H.C. Man, Cavitation erosion resistance of AISI 316L stainless steel laser surface-modified with NiTi, Mater. Sci. Eng. A 392 (2005) 348–358. [57] W. Liu, Y.G. Zheng, C.S. Liu, Z.M. Yao, Cavitation erosion characteristics of a NiTi alloy, Metall. Mater. Trans. A 35 (2004) 356–362. [58] J.B. Vogt, T. Magnin, J. Foct, Effective stresses and microstructure in cyclically deformed 316L austenitic stainless steel: effect of temperature and nitrogen content, Fatigue Fract. Eng. M 16 (2010) 555–564. [59] J. Polák, F. Fardoun, S. Degallaix, Analysis of the hysteresis loop in stainless steels II. austenitic-ferritic duplex steel and the effect of nitrogen, Mater. Sci. Eng. A 297 (2001) 154–161. [60] R. Fawley, M.A. Quader, R.A. Dodd, Compositional effects on the deformation modes, annealing twin frequencies and stacking fault energies of austenitic stainless steels, Trans. Metall. Soc. AIME 242 (1968) 771–776. [61] P.R. Swann, Dislocation substructure vs transgranular stress corrosion susceptibility of single phase alloys, Corrosion 19 (1963) 102–114. [62] R.E. Schramm, R.P. Reed, Stacking fault energies of seven commercial austenitic stainless steels, Metall. Trans. A 6 (1975) 1345–1351.
Microstructure, mechanical and corrosion properties of friction stir welded high nitrogen nickel-free austenitic stainless steel, Mater. Des. 84 (2015) 291–299. M. Moallemi, A. Zarei-Hanzaki, H.S. Baghbadorani, Evolution of microstructure and mechanical properties in a cold deformed nitrogen bearing TRIP-assisted duplex stainless steel after reversion annealing, Mater. Sci. Eng. A 683 (2017) 83–89. A. Di Schino, M.G. Mecozzi, M. Barteri, Solidification mode and residual ferrite in low-Ni austenitic stainless steels, J. Mater. Sci. 35 (2000) 375–380. H.B. Li, Z.H. Jiang, Y. Cao, Z.R. Zhang, Fabrication of high nitrogen austenitic stainless steels with excellent mechanical and pitting corrosion properties, Int. J. Miner. Metall. Mater. 16 (2009) 387–392. M. Talha, C.K. Behera, O.P. Sinha, Effect of nitrogen and cold working on structural and mechanical behavior of Ni-free nitrogen containing austenitic stainless steels for biomedical applications, Mater. Sci. Eng. C 47 (2015) 196–203. L.M. Kaputkina, V.G. Prokoshkina, Martensitic transformations and martensite structure in thermomechanically strengthened high-nitrogen steels, Mater. Sci. Eng. A 438–440 (2006) 228–232. Y.X. Qiao, Y.G. Zheng, W. Ke, P.C. Okafor, Electrochemical behaviour of high nitrogen stainless steel in acidic solutions, Corros. Sci. 51 (2009) 979–986. Y.X. Qiao, X. Cai, J. Cui, H.B. Li, Passivity and semiconducting behavior of a high nitrogen stainless steel in acidic NaCl solution, Adv. Mater. Sci. Eng. (2016) (No.6065481). P. Wan, Y.B. Ren, B.C. Zhang, K. Yang, Effect of nitrogen on biocorrosion behavior of high nitrogen Nickel-free stainless steel in different simulated body fluids, Mater. Sci. Eng. C 32 (2012) 510–516. Y.X. Qiao, Z.H. Tian, X. Cai, J. Chen, Y.X. Wang, Q.N. Song, H.B. Li, Cavitation erosion behaviors of a nickel-free high-nitrogen stainless steel, Tribol. Lett. 67 (2019) 1–9. W.T. Fu, Y.Z. Zheng, X.K. He, Resistance of a high nitrogen austenitic steel to cavitation erosion, Wear 249 (2001) 788–791. W. Liu, Y.G. Zheng, C.S. Liu, Z.M. Yao, W. Ke, Cavitation erosion behavior of Cr-MnN stainless steels in comparison with 0Cr13Ni5Mo stainless steel, Wear 254 (2003) 713–722. G. Bregliozzia, A. Di Schinob, S.I.-U. Ahmeda, J.M. Kenny, H. Haefkea, Cavitation wear behaviour of austenitic stainless steels with different grain sizes, Wear 258 (2005) 503–510. D.H. Mesa, C.M. Garzón, A.P. Tschiptschin, Influence of cold-work on the cavitation erosion resistance and on the damage mechanisms in high-nitrogen austenitic stainless steels, Wear 271 (2011) 1372–1377. P. Niederhofer, S. Huth, Cavitation erosion resistance of high interstitial CrMnCN austenitic stainless steels, Wear 301 (2013) 457–466. Y.X. Qiao, S. Wang, B. Liu, Y.G. Zheng, H.B. Li, Z.H. Jiang, Synergistic effect of corrosion and cavitation erosion of high nitrogen stainless steel, Acta Metall. Sin. 52 (2016) 233–240. J. Li, Y.X. Yang, Y.B. Ren, J.H. Dong, K. Yang, Effect of cold deformation on corrosion fatigue behavior of nickel-free high nitrogen austenitic stainless steel for coronary stent application, J. Mater. Sci. Technol. 34 (2018) 660–665. J.W. Simmons, Influence of nitride (Cr2N) precipitation on the plastic flow behavior of high-nitrogen austenitic stainless steel, Scr. Metall. 32 (1995) 265–270. H. Lee, S.J. Kim, Y.C. Jung, Crystallographic details of precipitates in Fe-22Cr-21Ni6Mo-(N) superaustenitic stainless steels aged at 900 °C, Metall. Mater. Trans. A 31 (2000) 1713–1723. R.D. Knutsen, C.L. Lang, J.A. Basson, Discontinuous cellular precipitation in a CrMn-N steel with niobium and vanadium additions, Acta Mater. 52 (2004) 2407–2417. H.B. Li, Z.H. Jiang, H. Feng, Q.F. M, D.P. Zhan, Aging precipitation behavior of 18Cr-16Mn-2Mo-1.1N high nitrogen austenitic stainless steel and its influences on mechanical properties, J. Iron Steel Res. Int. 19 (2012) 43–51. N. Sathirachinda, R. Pettersson, J. Pan, Depletion effects at phase boundaries in 2205 duplex stainless steel characterized with SKPFM and TEM/EDS, Corros. Sci. 51 (2009) 1850–1860. Y.H. Yang, B. Yan, J. Wang, J.L. Yin, The influence of solution treatment temperature on microstructure and corrosion behavior of high temperature ageing in 25% Cr duplex stainless steel, J. Alloy Compd. 509 (2011) 8870–8879. M.Q. Ou, Y.C. Ma, X.C. Hao, B.F. Wan, T. Liang, K. Liu, Effect of solution annealing on microstructure and mechanical properties of a Ni-Cr-W-Fe alloy, J. Mater. Sci. Technol. 33 (2017) 1300–1307. C.T. Kwok, H.C. Man, F.T. Cheng, Cavitation erosion and damage mechanisms of alloys with duplex structures, Mater. Sci. Eng. A 242 (1998) 108–120. Z.Y. Wang, J.H. Zhu, Correlation of martensitic transformation and surface mechanical behavior with cavitation erosion resistance for some iron-based alloys, Wear 256 (2004) 1208–1213. R.B. Nair, H.S. Arora, S. Mukherjee, S. Singh, H. Singh, H.S. Grewal, Exceptionally high cavitation erosion and corrosion resistance of a high entropy alloy, Ultrason. Sonochem. 41 (2018) 252–260. F.T. Cheng, P. Shi, H.C. Man, Correlation of cavitation erosion resistance with indentation-derived properties for a NiYi alloy, Scr. Mater. 45 (2001) 1083–1089. I. Mitelea, M.L. Micu, I. Bordeasu, C.M. Crãciunescu, Cavitation erosion of sensitized UNS S31803 duplex stainless steels, J. Mater. Eng. Perform. 25 (2016) 1939–1944.
77