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Microstructural characterization of low temperature plasma-nitrided 316L stainless steel surface with prior severe shot peening
Materials and Design 108 (2016) 448–454
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Microstructural characterization of low temperature plasma-nitrided 316L stainless steel surface with prior severe shot peening M. Jayalakshmi a,⁎, Prashant Huilgol a, B. Ramachandra Bhat b, K. Udaya Bhat a,⁎ a b
Dept. of Metallurgical and Materials Engineering, National Institute of Technology Karnataka, Surathkal, Karnataka575025, India Dept. of Chemistry, National Institute of Technology Karnataka, Surathkal, Karnataka575025, India
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• Severe shot peening is successfully used as a pre-treatment for low temperature plasma nitriding of 316L stainless steels. • Significantly high, uniform case depth of 45 μm is achieved by plasma nitriding at 400 °C for 4 hours. • Precipitation of deleterious chromium nitride is not observed. • Major phase in the nitride layer is observed to be martensitic in nature.
a r t i c l e
i n f o
Article history: Received 29 April 2016 Received in revised form 13 June 2016 Accepted 3 July 2016 Available online 04 July 2016 Keywords: Surface nanocrystallization Severe shot-peening Plasma nitriding Austenitic stainless steel Martensite
a b s t r a c t Surface nanocrystallization by severe deformation has proven beneficial as pre-treatment to plasma nitriding. It aids in achieving thicker nitride layers at lower temperatures thus making the process more economical. In austenitic stainless steels, severe deformation leads to formation of strain induced martensite on the surface while plasma nitriding alone forms expanded austenite. However, structural characteristics of surface layer of pre-deformed steel after plasma nitriding is still a matter of debate. In present study, 316L stainless steel was subjected to severe shot peening: followed by plasma nitriding at 400 °C for 4 h. Characteristics of sample surface before and after treatment were analyzed by scanning electron microscopy, X-ray diffractometry and transmission electron microscopy techniques. Results showed that, this duplex treatment leads to formation of about 45 μm thick nitride layer; without CrN precipitation. This is significantly high compared to reported data considering the temperature and duration of nitriding treatment employed. Selected area electron diffraction pattern from topmost surface confirmed the co-existence of austenite and martensite while subsurface layer was predominantly consisting of lath martensite. This indicates that major phase in the nitrided layer is martensitic in nature and nitrogen supersaturation leads to transformation of small fraction of martensite to expanded austenite. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction ⁎ Corresponding author. E-mail addresses: [email protected] (M. Jayalakshmi), [email protected] (K.U. Bhat).
http://dx.doi.org/10.1016/j.matdes.2016.07.005 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
Low-temperature plasma nitriding is a well established surface engineering technique for producing hard and wear resistant surface layers on the austenitic stainless steels [1–4]. Nitriding in the temperature
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range of 400–450 °C prevents the formation of detrimental CrN [5–7]; but the case depth achieved would be significantly less owing to slower diffusion of nitrogen through austenite lattice. It is already proven that nanocrystalline and ultra-fine grains on the substrate surface along with higher defect concentration aids in faster diffusion [8]. Hence, several attempts were made to utilize severe plastic deformation techniques as pre-treatment for low temperature plasma nitriding. Surface nanocrystallization was achieved by surface mechanical attrition treatment (SMAT) [9–12], shot peening [13–16], high pressure torsion (HPT) [17], cold rolling [18], etc. Plasma nitriding following severe plastic deformation yielded appreciably thick nitride layers. Generally, severe plastic deformation leads to transformation of austenite phase to strain induced martensite [19], while, low temperature plasma nitriding alone leads to the formation of expanded austenite (known as S-phase) in the austenitic stainless steels [20–22]. But when plasma nitriding of pre-deformed steels is carried out, several studies [12,15,23–25] advocate the complete transformation of strain induced martensite to Sphase over nitriding while studies by Ferkel et al. [17] and Ji et al. [14] suggested that the phases formed in the nitride layer could be a mixture of ferritic and/or martensitic phases along with Fe and/or Cr nitrides. Hence, further experiments are necessary to elucidate microstructure of the surface layer after the duplex treatment. In the present study, microstructural evolution of 316L stainless steel after severe shot peening followed by plasma nitriding treatment was studied to understand the phase/s present in the nitride layer. 2. Experimental details 316L stainless steel (SS) sheets [Composition: Cr 17.8%, Ni 11.5%, C 0.07%, Mn 1.58%, Si 0.04%, Ti 0.64%, S 0.006%, P 0.012%, Fe bal., all in wt%] of 5 mm thickness in hot rolled condition were used in the study. 10 cm ∗ 10 cm coupons of steel were cut, polished metallographically to a roughness of 0.2 μm and cleaned ultrasonically in acetone and distilled water. Cleaned samples were subjected to air blast shot peening (M/s Curtiss Wright surface technologies, Bangalore). Standard cast steel shots, S230, having 0.6 mm diameter were used for peening and shot flow was maintained at 25 kg/min. Shot peening intensity as measured on “Almen A” strip was 8 and samples were peened with 1000% coverage to ensure severe peening. Pre-peened and un-peened (as received) samples were subjected to plasma nitriding (M/s Bhat Metals, Pune) at 400 °C for 4 h with a N2 to H2 ratio of 1:1. As received, peened and peened-nitrided samples were metallographically polished and etched in dilute aqua regia (HCl:HNO3:water is 2:1:1) to study the microstructural features using scanning electron microscope (SEM, 6380 LA, JEOL). Phases were identified through X-Ray diffractometer (XRD, JDX 8P, JEOL) using a monochromatic CuKα radiation of wavelength 0.154 nm. Data was collected in diffracting angle range of 30 to 90° at a step size of 0.02° and scan speed of 1° per minute. Transmission electron microscope (TEM, JEM-2100, JEOL) operating at 200 kV was used to study the microstructural constituents in the surface and subsurface layers. Samples for transmission electron microscopy studies were prepared by thinning from the un-treated side to the site of interest; followed by punching, dimpling and ion milling. Ion milling was done using Gatan PIPS instrument at an accelerating voltage of 5 kV and angle of 5°. 3. Results 3.1. SEM observations The typical microstructure of the as- received 316L SS sample is shown in Fig. 1. Austenite grains of size in a range of 40–80 μm are visible; with annealing twins in few grains. Cross sectional SEM micrographs of the sample after severe shot peening are shown in Fig. 2. Based on the microstructural features, three regions are identified in Fig. 2(a). They are (from the top): region
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Fig. 1. SEM micrograph of the as-received AISI 316L stainless steel (Etchant-dilute aqua regia).
A which has undergone nanocrystallization, second is the region B which is characterized by the presence of planar dislocation arrays, mechanical twins, rhombic blocks produced by multidirection twin intersections, etc. and last is the unaffected base material. Higher magnification image from the region A in Fig. 2(b) indicates that coarse austenite grains of base material formed during hot-rolling have undergone severe deformation to form finer grains. Formation of nanocrystalline grains are confirmed through TEM studies. Fig. 2(c) is a typical micrograph in the subsurface region B. It shows twin-twin intersections resulting in the formation of rhombic blocks [26,27] of 1-5 μm thickness. Cross-sectional SEM micrographs of peened and un-peened samples after plasma nitriding are shown in Fig. 3. It can be observed that the nitride layer obtained in case of un-peened sample is less than 1 μm in thickness while severely peened sample possesses a well defined, uniform nitride layer of about 45 μm thickness. Also, nitride layer appears featureless with no cracks or precipitates while the base material is attacked by the etchant. A gradual transition from nitride layer to unaffected base material is observed through a transition zone consisting of deformation features induced during shot peening.
3.2. XRD results The X-ray diffraction patterns of as-received, peened and peened-nitrided samples are shown in Fig. 4. All peaks of as- received 316L SS are corresponding to austenite phase. After peening, XRD pattern predominantly consists of martensite phase with a small fraction of austenite. This confirms the strain induced transformation of majority of austenite to martensite due to severe deformation. It is interesting to note that after plasma nitriding of pre-peened sample, a broad peak in the 2θ range of 40–43° and a sharp peak at diffraction angle of 44.7° co-exist in the XRD pattern. The broad peak could be assigned to the (111) plane of expanded austenite which is observed in several investigations [20,28,29] on plasma nitriding of austenitic stainless steels. The (200) reflection from CrN and (110) from martensite phase are probable in the diffraction angle of about 44°. Hence several studies have attributed the observed sharp peak to the precipitation of CrN [10,13,15]. However, few authors have left the peak unidentified stating that it could be a mixture of ferritic and/or martensitic phases along with Fe and/or Cr nitrides [14,17]. Hence TEM studies were done to confirm the phases present in the treated layer. However, observed peak at 78.2° was not matching to any of the phases (austenite: JCPDS card no. 33-0397, martensite: JCPDS card no. 35-1375, CrN: JCPDS card no. 11-0065, expanded austenite) in the nitride layer and is left unidentified.
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Fig. 2. (a) Cross-sectional SEM micrograph of 316L SS after severe shot peening, typical magnified image from (b) region A and (c) region B.
3.3. TEM studies 3.3.1. Surface layer after severe shot peening TEM micrograph and corresponding selected area electron diffraction (SAED) pattern from the top most layer of the severely peened sample is given in Fig. 5. Nanocrystalline grains of martensite with random crystallographic orientation are observed. No austenite rings are present in any of the SAED patterns taken from several places in the sample; indicating complete transformation at the surface. However, presence of small fraction of austenite cannot be neglected as indicated by the XRD data in Fig.4; as the area analyzed by SAED patterns will be relatively less. The martensite formed is of dislocation cell-type; which is the characteristic of austenite transformed to martensite at higher strain and strain rate [30]. 3.3.2. Surface layer after plasma nitriding Fig. 6 shows TEM micrograph and corresponding SAED patterns of top most layer of peened-nitrided sample. Grains in the range of 30– 100 nm are seen in Fig. 6(a) suggesting that the ultrafine grain structure obtained during shot peening is retained even after nitriding. However, dislocation density within the grains was notably less. SAED pattern in Fig. 6(b) indicates the presence of FCC diffraction rings of along with those of martensite. According to Yu et al., nitrogen being an austenite stabilizer, supersaturation of nitrogen atoms in the lattice of martensite
triggers the transformation of martensite to expanded austenite. XRD pattern of nitride sample in Fig. 4 also confirms the presence of expanded austenite. Hence, the FCC rings may be attributed to expanded austenite phase. It is also probable that the expanded austenite is formed by the incorporation of nitrogen atoms into the lattice of untransformed parent austenite phase during plasma nitriding. TEM micrograph and corresponding SAED pattern from 15 μm below (approximately) the top layer of peened-nitrided sample are given in Fig. 7. Presence of lath martensite is clearly visible in Fig. 7(a) with individual lath having thickness in the range of 60–100 nm. Each lath is found to consist of huge amount of dislocation sub-structures. Corresponding SAED pattern in Fig. 7(b) confirms the presence of only martensite phase. Precipitation of CrN was not observed either on the top surface or in the sub-surface layer. Hence the peak corresponding to 2θ of 44.7° in the XRD pattern of peened-nitrided sample in Fig. 3 can be undoubtedly attributed to the martensite. 4. Discussion 4.1. Microstructure after severe shot peening From SEM micrographs (Fig. 2), XRD plots (Fig. 4) and TEM (Fig. 5) micrographs, it is evident that the surface of the steel has undergone transformation to produce martensite during shot peening. The
Fig. 3. Cross sectional SEM micrograph of (a) un-peened and (b) severely peened sample after plasma nitriding.
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4.2. Microstructure after plasma nitriding
Fig. 4. XRD patterns of unpeened, peened and peened-nitrided 316L.
martensite is very fine (nanocrystalline) and the units are oriented randomly. A schematic diagram of the surface nanocrystallization and martensite formation upon shot peening is shown in Fig. 8. During severe shot peening, the strain and strain rate experienced at the surface is markedly high. Strain induced by repeated impacts on the sample surface favors the formation of planar dislocation arrays and mechanical twins (Fig. 8a). In low stacking fault energy materials like austenitic stainless steels, twinning is the dominant deformation mechanism specifically at higher strain rates and/or lower temperatures [31]. With increase in strain, the intersection of multidirectional twins takes place leading to formation of rhombic blocks (Fig. 8b). These intersections repeatedly subdivide the original coarse austenite grains and subsequently lead to nanocrystallization [26,27,32]. It has been shown that strain induced martensite nucleates at twin intersections (Fig. 8c) [33,34] and grow further to facilitate complete transformation of austenite to martensite. Hence, on the top surface of peened sample, nano-sized grains are observed as shown in Fig. 2(b) and complete transformation of austenite to strain induced martensite is evident in Fig. 5(b). Typically, strain induced martensite in austenitic stainless steels grows in lath morphology [35]. However, it is shown that when the strain and/or strain rate encountered during the deformation are high, lath structure is broken, refinement of lath packets takes place; thus leading to the formation of dislocation cell-type martensite [30]. This explains the observed blocky structure of martensite in Fig. 5(a) involving a large amount of dislocations.
Fig. 6(a) shows a TEM micrograph at the top surface. Comparing it with Fig. 5(a), it is clear that the dislocation density is reduced during plasma nitriding treatment, wherein the samples were maintained at 400 °C for 4 h. It is established that the thermal energy assists in recombination, rearrangement and annihilation of dislocations substructures [36]. This could be the reason for the observed decrease in the dislocation density in Fig. 6(a) on sample surface, after plasma nitriding. It may be noted that the morphology of the martensite of the topmost layer remains same even after plasma nitriding. It is interesting to note that, for the nitriding temperature and duration employed in the present study, the nitride case depth obtained is appreciably high compared to previous studies [13–15]. Shot peening was used as pre-treatment in all the studies. Though nitride layer thickness of 50 μm is reported in [13], gas nitriding was carried out at 570 °C for 8 h. In the work of Ji et al. [14] nitride layer thickness achieved was about 20 μm inspite of nitriding duration of at 400 °C for 6 h. Case depth obtained was limited to 5 μm in the work of Shen et al. [15] when plasma nitriding was carried out at 400 °C for 4 h. The case depth of 45 μm achieved in the present study is attributed to enhanced nitriding kinetics due to the surface nanocrystallization and presence of strain induced martensite. Higher free energy state near the surface, higher volume fraction of the grain boundaries, triple junctions, twin boundaries and high dislocation densities provide easy path for the diffusion of nitrogen into the sample. Also, the diffusion coefficient of nitrogen in martensite is higher than that in austenite [25]. Hence higher fraction of martensite facilitates rapid diffusion of nitrogen. It is well accepted that low temperature plasma nitriding of austenitic stainless steel results in the formation of expanded austenite in nitride layer. In the last decade, severe plastic deformation technique has emerged as a successful pre-treatment step for enhancement of nitriding kinetics in the austenitic stainless steels. But till date, there is no agreement in views amongst the researchers about the microstuctural evolution and phases formed during this duplex process. In some studies, no martensite is reported after SMAT pre-treatment [10,11], while austenite transforms to S-phase after plasma nitriding. Few authors opine that the strain induced martensite formed during the deformation gets converted into expanded austenite on plasma nitriding; which further transforms into α-Fe and CrN upon nitriding at higher temperatures for longer durations [9,37]. In few cases, Fe4N phase is also reported [13]. In the present investigation, there is no indication of CrN, suggesting that corrosion resistance of the steel is not impaired
Fig. 5. (a) TEM micrograph and (b) SAED pattern of severely shot peened sample.
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Fig. 6. (a) TEM micrograph and (b) corresponding SAED pattern from top surface of peened-nitrided sample.
as chromium has not diffused out from the matrix. This is distinctly different than investigations [9,37] in which extensive formation of CrN during nitriding is reported. Yu et al. [25] proposed a two- step transformation mechanism from martensite to expanded austenite. It suggests a complete transformation of strain induced martensite to austenite initially, followed by the transformation of austenite to S-phase. Several studies [12,15,24] have justified their results based on this theory. Further, a study by Ferkel et al. [17] indicates the presence of martensite in nitride layer. It suggests that the martensite could be due to the strain induced transformation of austenite during prior severe deformation or due to the decomposition of expanded austenite to martensite and CrN. The study by Ji et al. [14] confirmed the formation of strain induced martensite upon deformation, while phases formed after plasma nitriding were left unidentified suggesting that it could be a mixture of ferritic and/or martensitic phases along with Fe and/or Cr nitrides. This ambiguity was due to observed broad XRD peaks in the region of 40–45° and hence, TEM analysis is required to confirm the phases present. In a recent study [23] also, it has been stated that the presence of strain induced martensite enhances the nitriding layer thickness and it transforms to expanded austenite when nitrided at lower temperatures while it promotes CrN formation at higher treatment temperatures. In the present study it is confirmed that, under the process parameters
used for the duplex treatment, precipitation of deleterious CrN is not observed. It is also evident from the TEM studies that the major phase in the surface and subsurface layer is martensite. Hence, it is reasonable to assume that the nitrogen has entered into martensite lattice during nitriding treatment to facilitate the formation of about 45 μm thick nitride layer. However, no peak shift or broadening is observed in the XRD peaks of martensite due to nitrogen incorporation. Nitrogen being an austenite stabilizer, its supersaturation at the surface facilitates the transformation of martensite to expanded austenite in few regions as reported by Yu et al. [25]. It is evident by appearance of broad peak of S-phase in Fig.4 and weak diffraction rings of FCC-structure in Fig. 6(b). This study differs from most of the previous reports indicating that the major phase in the nitride layer is martensitic in nature. 5. Conclusion The study shows primarily that a nitride case depth of about 45 μm is successfully produced by plasma nitriding of 316L stainless steel at 400 ° C with shot peening pre-treatment. Nitriding case depth obtained is much higher compared to the reported literature for nitriding duration of 4 h. Shot-peening as a pre-treatment step induced nanocrystalline surface layers on the sample having martensite formed due to strain
Fig. 7. (a) TEM micrograph and (b) corresponding SAED pattern of peened-nitrided sample from about 15 μm below the surface.
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Fig. 8. Schematic diagram of formation of martensite during shot peening.
induced transformation of austenite. Nanocrystallinity is retained even after plasma nitriding treatment. Top layer of nitride is found to consist of majority of martensitic phase; while sub-surface (about 15 μm below the surface) found to possess lath martensite structure. Weak reflections of S-phase in XRD and diffraction rings corresponding to FCC structure is attributed to the transformation of a fraction martensite to Sphase, due to supersaturation of nitrogen in the martensite lattice. Hence it is inferred that the nitride layer is formed largely by incorporation of nitrogen into martensite lattice; contrary to majority of the reports. Also, CrN is completely absent due to the lower nitriding temperature employed in the study. Yet, thick nitride layer is obtained, which is attributed to extensive formation of defects and martensite during shot-peening pre-treatment. Acknowledgements We would like to acknowledge the Director, NITK and MHRD, Govt. of India for providing financial assistance for this research work. References [1] Y. Sun, T. Bell, Sliding wear characteristics of low temperature plasma nitrided 316 austenitic stainless steel, Wear 218 (1998) 34–42, http://dx.doi.org/10.1016/ S0043-1648(98)00199-9. [2] E. Menthe, A. Bulak, J. Olfe, A. Zimmermann, K.T. Rie, Improvement of the mechanical properties of austenitic stainless steel after plasma nitriding, Surf. Coat. Technol. 133–134 (2000) 259–263, http://dx.doi.org/10.1016/S0257-8972(00)00930-0. [3] E. Menthe, K.T. Rie, Further investigation of the structure and properties of austenitic stainless steel after plasma nitriding, Surf. Coat. Technol. 116-119 (1999) 199–204, http://dx.doi.org/10.1016/S0257-8972(99)00085-7. [4] W. Liang, Surface modification of AISI 304 austenitic stainless steel by plasma nitriding, Appl. Surf. Sci. 211 (2003) 308–314, http://dx.doi.org/10.1016/S01694332(03)00260-5. [5] Z.L. Zhang, T. Bell, Structure and corrosion resistance of plasma nitrided stainless steel, Surf. Eng. 1 (1985) 131–136, http://dx.doi.org/10.1179/sur.1985.1.2.131. [6] C. Blawert, A. Weisheit, B.L. Mordike, R.M. Knoop, Plasma immersion ion implantation of stainless steel: austenitic stainless steel in comparison to austenitic-ferritic stainless steel, Surf. Coat. Technol. 85 (1996) 15–27, http://dx.doi.org/10.1016/ 0257-8972(96)02880-0. [7] L. Gil, S. Brühl, L. Jiménez, O. Leon, R. Guevara, M.H. Staia, Corrosion performance of the plasma nitrided 316L stainless steel, Surf. Coat. Technol. 201 (2006) 4424–4429, http://dx.doi.org/10.1016/j.surfcoat.2006.08.081. [8] W.P. Tong, Nitriding iron at lower temperatures, Science 299 (2003) 686–688, http://dx.doi.org/10.1126/science.1080216 (80-.). [9] Y. Lin, J. Lu, L. Wang, T. Xu, Q. Xue, Surface nanocrystallization by surface mechanical attrition treatment and its effect on structure and properties of plasma nitrided AISI 321 stainless steel, Acta Mater. 54 (2006) 5599–5605, http://dx.doi.org/10.1016/j. actamat.2006.08.014. [10] M. Laleh, F. Kargar, M. Velashjerdi, Low-temperature nitriding of nanocrystalline stainless steel and its effect on improving wear and corrosion resistance, J. Mater. Eng. Perform. 22 (2013) 1304–1310, http://dx.doi.org/10.1007/s11665-012-0417-7. [11] M. Chemkhi, D. Retraint, A. Roos, C. Garnier, L. Waltz, C. Demangel, et al., The effect of surface mechanical attrition treatment on low temperature plasma nitriding of an austenitic stainless steel, Surf. Coat. Technol. 221 (2013) 191–195, http://dx.doi.org/ 10.1016/j.surfcoat.2013.01.047. [12] T. Balusamy, T.S.N.S. Narayanan, K. Ravichandran, I.S. Park, M.H. Lee, Plasma nitriding of AISI 304 stainless steel: role of surface mechanical attrition treatment, Mater. Charact. 85 (2013) 38–47, http://dx.doi.org/10.1016/j.matchar.2013.08.009.
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Microstructural characterization of low temperature plasma-nitrided 316L stainless steel surface with prior severe shot peening M. Jayalakshmi a,⁎, Prashant Huilgol a, B. Ramachandra Bhat b, K. Udaya Bhat a,⁎ a b
Dept. of Metallurgical and Materials Engineering, National Institute of Technology Karnataka, Surathkal, Karnataka575025, India Dept. of Chemistry, National Institute of Technology Karnataka, Surathkal, Karnataka575025, India
H I G H L I G H T S
G R A P H I C A L
A B S T R A C T
• Severe shot peening is successfully used as a pre-treatment for low temperature plasma nitriding of 316L stainless steels. • Significantly high, uniform case depth of 45 μm is achieved by plasma nitriding at 400 °C for 4 hours. • Precipitation of deleterious chromium nitride is not observed. • Major phase in the nitride layer is observed to be martensitic in nature.
a r t i c l e
i n f o
Article history: Received 29 April 2016 Received in revised form 13 June 2016 Accepted 3 July 2016 Available online 04 July 2016 Keywords: Surface nanocrystallization Severe shot-peening Plasma nitriding Austenitic stainless steel Martensite
a b s t r a c t Surface nanocrystallization by severe deformation has proven beneficial as pre-treatment to plasma nitriding. It aids in achieving thicker nitride layers at lower temperatures thus making the process more economical. In austenitic stainless steels, severe deformation leads to formation of strain induced martensite on the surface while plasma nitriding alone forms expanded austenite. However, structural characteristics of surface layer of pre-deformed steel after plasma nitriding is still a matter of debate. In present study, 316L stainless steel was subjected to severe shot peening: followed by plasma nitriding at 400 °C for 4 h. Characteristics of sample surface before and after treatment were analyzed by scanning electron microscopy, X-ray diffractometry and transmission electron microscopy techniques. Results showed that, this duplex treatment leads to formation of about 45 μm thick nitride layer; without CrN precipitation. This is significantly high compared to reported data considering the temperature and duration of nitriding treatment employed. Selected area electron diffraction pattern from topmost surface confirmed the co-existence of austenite and martensite while subsurface layer was predominantly consisting of lath martensite. This indicates that major phase in the nitrided layer is martensitic in nature and nitrogen supersaturation leads to transformation of small fraction of martensite to expanded austenite. © 2016 Elsevier Ltd. All rights reserved.
1. Introduction ⁎ Corresponding author. E-mail addresses: [email protected] (M. Jayalakshmi), [email protected] (K.U. Bhat).
http://dx.doi.org/10.1016/j.matdes.2016.07.005 0264-1275/© 2016 Elsevier Ltd. All rights reserved.
Low-temperature plasma nitriding is a well established surface engineering technique for producing hard and wear resistant surface layers on the austenitic stainless steels [1–4]. Nitriding in the temperature
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range of 400–450 °C prevents the formation of detrimental CrN [5–7]; but the case depth achieved would be significantly less owing to slower diffusion of nitrogen through austenite lattice. It is already proven that nanocrystalline and ultra-fine grains on the substrate surface along with higher defect concentration aids in faster diffusion [8]. Hence, several attempts were made to utilize severe plastic deformation techniques as pre-treatment for low temperature plasma nitriding. Surface nanocrystallization was achieved by surface mechanical attrition treatment (SMAT) [9–12], shot peening [13–16], high pressure torsion (HPT) [17], cold rolling [18], etc. Plasma nitriding following severe plastic deformation yielded appreciably thick nitride layers. Generally, severe plastic deformation leads to transformation of austenite phase to strain induced martensite [19], while, low temperature plasma nitriding alone leads to the formation of expanded austenite (known as S-phase) in the austenitic stainless steels [20–22]. But when plasma nitriding of pre-deformed steels is carried out, several studies [12,15,23–25] advocate the complete transformation of strain induced martensite to Sphase over nitriding while studies by Ferkel et al. [17] and Ji et al. [14] suggested that the phases formed in the nitride layer could be a mixture of ferritic and/or martensitic phases along with Fe and/or Cr nitrides. Hence, further experiments are necessary to elucidate microstructure of the surface layer after the duplex treatment. In the present study, microstructural evolution of 316L stainless steel after severe shot peening followed by plasma nitriding treatment was studied to understand the phase/s present in the nitride layer. 2. Experimental details 316L stainless steel (SS) sheets [Composition: Cr 17.8%, Ni 11.5%, C 0.07%, Mn 1.58%, Si 0.04%, Ti 0.64%, S 0.006%, P 0.012%, Fe bal., all in wt%] of 5 mm thickness in hot rolled condition were used in the study. 10 cm ∗ 10 cm coupons of steel were cut, polished metallographically to a roughness of 0.2 μm and cleaned ultrasonically in acetone and distilled water. Cleaned samples were subjected to air blast shot peening (M/s Curtiss Wright surface technologies, Bangalore). Standard cast steel shots, S230, having 0.6 mm diameter were used for peening and shot flow was maintained at 25 kg/min. Shot peening intensity as measured on “Almen A” strip was 8 and samples were peened with 1000% coverage to ensure severe peening. Pre-peened and un-peened (as received) samples were subjected to plasma nitriding (M/s Bhat Metals, Pune) at 400 °C for 4 h with a N2 to H2 ratio of 1:1. As received, peened and peened-nitrided samples were metallographically polished and etched in dilute aqua regia (HCl:HNO3:water is 2:1:1) to study the microstructural features using scanning electron microscope (SEM, 6380 LA, JEOL). Phases were identified through X-Ray diffractometer (XRD, JDX 8P, JEOL) using a monochromatic CuKα radiation of wavelength 0.154 nm. Data was collected in diffracting angle range of 30 to 90° at a step size of 0.02° and scan speed of 1° per minute. Transmission electron microscope (TEM, JEM-2100, JEOL) operating at 200 kV was used to study the microstructural constituents in the surface and subsurface layers. Samples for transmission electron microscopy studies were prepared by thinning from the un-treated side to the site of interest; followed by punching, dimpling and ion milling. Ion milling was done using Gatan PIPS instrument at an accelerating voltage of 5 kV and angle of 5°. 3. Results 3.1. SEM observations The typical microstructure of the as- received 316L SS sample is shown in Fig. 1. Austenite grains of size in a range of 40–80 μm are visible; with annealing twins in few grains. Cross sectional SEM micrographs of the sample after severe shot peening are shown in Fig. 2. Based on the microstructural features, three regions are identified in Fig. 2(a). They are (from the top): region
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Fig. 1. SEM micrograph of the as-received AISI 316L stainless steel (Etchant-dilute aqua regia).
A which has undergone nanocrystallization, second is the region B which is characterized by the presence of planar dislocation arrays, mechanical twins, rhombic blocks produced by multidirection twin intersections, etc. and last is the unaffected base material. Higher magnification image from the region A in Fig. 2(b) indicates that coarse austenite grains of base material formed during hot-rolling have undergone severe deformation to form finer grains. Formation of nanocrystalline grains are confirmed through TEM studies. Fig. 2(c) is a typical micrograph in the subsurface region B. It shows twin-twin intersections resulting in the formation of rhombic blocks [26,27] of 1-5 μm thickness. Cross-sectional SEM micrographs of peened and un-peened samples after plasma nitriding are shown in Fig. 3. It can be observed that the nitride layer obtained in case of un-peened sample is less than 1 μm in thickness while severely peened sample possesses a well defined, uniform nitride layer of about 45 μm thickness. Also, nitride layer appears featureless with no cracks or precipitates while the base material is attacked by the etchant. A gradual transition from nitride layer to unaffected base material is observed through a transition zone consisting of deformation features induced during shot peening.
3.2. XRD results The X-ray diffraction patterns of as-received, peened and peened-nitrided samples are shown in Fig. 4. All peaks of as- received 316L SS are corresponding to austenite phase. After peening, XRD pattern predominantly consists of martensite phase with a small fraction of austenite. This confirms the strain induced transformation of majority of austenite to martensite due to severe deformation. It is interesting to note that after plasma nitriding of pre-peened sample, a broad peak in the 2θ range of 40–43° and a sharp peak at diffraction angle of 44.7° co-exist in the XRD pattern. The broad peak could be assigned to the (111) plane of expanded austenite which is observed in several investigations [20,28,29] on plasma nitriding of austenitic stainless steels. The (200) reflection from CrN and (110) from martensite phase are probable in the diffraction angle of about 44°. Hence several studies have attributed the observed sharp peak to the precipitation of CrN [10,13,15]. However, few authors have left the peak unidentified stating that it could be a mixture of ferritic and/or martensitic phases along with Fe and/or Cr nitrides [14,17]. Hence TEM studies were done to confirm the phases present in the treated layer. However, observed peak at 78.2° was not matching to any of the phases (austenite: JCPDS card no. 33-0397, martensite: JCPDS card no. 35-1375, CrN: JCPDS card no. 11-0065, expanded austenite) in the nitride layer and is left unidentified.
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Fig. 2. (a) Cross-sectional SEM micrograph of 316L SS after severe shot peening, typical magnified image from (b) region A and (c) region B.
3.3. TEM studies 3.3.1. Surface layer after severe shot peening TEM micrograph and corresponding selected area electron diffraction (SAED) pattern from the top most layer of the severely peened sample is given in Fig. 5. Nanocrystalline grains of martensite with random crystallographic orientation are observed. No austenite rings are present in any of the SAED patterns taken from several places in the sample; indicating complete transformation at the surface. However, presence of small fraction of austenite cannot be neglected as indicated by the XRD data in Fig.4; as the area analyzed by SAED patterns will be relatively less. The martensite formed is of dislocation cell-type; which is the characteristic of austenite transformed to martensite at higher strain and strain rate [30]. 3.3.2. Surface layer after plasma nitriding Fig. 6 shows TEM micrograph and corresponding SAED patterns of top most layer of peened-nitrided sample. Grains in the range of 30– 100 nm are seen in Fig. 6(a) suggesting that the ultrafine grain structure obtained during shot peening is retained even after nitriding. However, dislocation density within the grains was notably less. SAED pattern in Fig. 6(b) indicates the presence of FCC diffraction rings of along with those of martensite. According to Yu et al., nitrogen being an austenite stabilizer, supersaturation of nitrogen atoms in the lattice of martensite
triggers the transformation of martensite to expanded austenite. XRD pattern of nitride sample in Fig. 4 also confirms the presence of expanded austenite. Hence, the FCC rings may be attributed to expanded austenite phase. It is also probable that the expanded austenite is formed by the incorporation of nitrogen atoms into the lattice of untransformed parent austenite phase during plasma nitriding. TEM micrograph and corresponding SAED pattern from 15 μm below (approximately) the top layer of peened-nitrided sample are given in Fig. 7. Presence of lath martensite is clearly visible in Fig. 7(a) with individual lath having thickness in the range of 60–100 nm. Each lath is found to consist of huge amount of dislocation sub-structures. Corresponding SAED pattern in Fig. 7(b) confirms the presence of only martensite phase. Precipitation of CrN was not observed either on the top surface or in the sub-surface layer. Hence the peak corresponding to 2θ of 44.7° in the XRD pattern of peened-nitrided sample in Fig. 3 can be undoubtedly attributed to the martensite. 4. Discussion 4.1. Microstructure after severe shot peening From SEM micrographs (Fig. 2), XRD plots (Fig. 4) and TEM (Fig. 5) micrographs, it is evident that the surface of the steel has undergone transformation to produce martensite during shot peening. The
Fig. 3. Cross sectional SEM micrograph of (a) un-peened and (b) severely peened sample after plasma nitriding.
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4.2. Microstructure after plasma nitriding
Fig. 4. XRD patterns of unpeened, peened and peened-nitrided 316L.
martensite is very fine (nanocrystalline) and the units are oriented randomly. A schematic diagram of the surface nanocrystallization and martensite formation upon shot peening is shown in Fig. 8. During severe shot peening, the strain and strain rate experienced at the surface is markedly high. Strain induced by repeated impacts on the sample surface favors the formation of planar dislocation arrays and mechanical twins (Fig. 8a). In low stacking fault energy materials like austenitic stainless steels, twinning is the dominant deformation mechanism specifically at higher strain rates and/or lower temperatures [31]. With increase in strain, the intersection of multidirectional twins takes place leading to formation of rhombic blocks (Fig. 8b). These intersections repeatedly subdivide the original coarse austenite grains and subsequently lead to nanocrystallization [26,27,32]. It has been shown that strain induced martensite nucleates at twin intersections (Fig. 8c) [33,34] and grow further to facilitate complete transformation of austenite to martensite. Hence, on the top surface of peened sample, nano-sized grains are observed as shown in Fig. 2(b) and complete transformation of austenite to strain induced martensite is evident in Fig. 5(b). Typically, strain induced martensite in austenitic stainless steels grows in lath morphology [35]. However, it is shown that when the strain and/or strain rate encountered during the deformation are high, lath structure is broken, refinement of lath packets takes place; thus leading to the formation of dislocation cell-type martensite [30]. This explains the observed blocky structure of martensite in Fig. 5(a) involving a large amount of dislocations.
Fig. 6(a) shows a TEM micrograph at the top surface. Comparing it with Fig. 5(a), it is clear that the dislocation density is reduced during plasma nitriding treatment, wherein the samples were maintained at 400 °C for 4 h. It is established that the thermal energy assists in recombination, rearrangement and annihilation of dislocations substructures [36]. This could be the reason for the observed decrease in the dislocation density in Fig. 6(a) on sample surface, after plasma nitriding. It may be noted that the morphology of the martensite of the topmost layer remains same even after plasma nitriding. It is interesting to note that, for the nitriding temperature and duration employed in the present study, the nitride case depth obtained is appreciably high compared to previous studies [13–15]. Shot peening was used as pre-treatment in all the studies. Though nitride layer thickness of 50 μm is reported in [13], gas nitriding was carried out at 570 °C for 8 h. In the work of Ji et al. [14] nitride layer thickness achieved was about 20 μm inspite of nitriding duration of at 400 °C for 6 h. Case depth obtained was limited to 5 μm in the work of Shen et al. [15] when plasma nitriding was carried out at 400 °C for 4 h. The case depth of 45 μm achieved in the present study is attributed to enhanced nitriding kinetics due to the surface nanocrystallization and presence of strain induced martensite. Higher free energy state near the surface, higher volume fraction of the grain boundaries, triple junctions, twin boundaries and high dislocation densities provide easy path for the diffusion of nitrogen into the sample. Also, the diffusion coefficient of nitrogen in martensite is higher than that in austenite [25]. Hence higher fraction of martensite facilitates rapid diffusion of nitrogen. It is well accepted that low temperature plasma nitriding of austenitic stainless steel results in the formation of expanded austenite in nitride layer. In the last decade, severe plastic deformation technique has emerged as a successful pre-treatment step for enhancement of nitriding kinetics in the austenitic stainless steels. But till date, there is no agreement in views amongst the researchers about the microstuctural evolution and phases formed during this duplex process. In some studies, no martensite is reported after SMAT pre-treatment [10,11], while austenite transforms to S-phase after plasma nitriding. Few authors opine that the strain induced martensite formed during the deformation gets converted into expanded austenite on plasma nitriding; which further transforms into α-Fe and CrN upon nitriding at higher temperatures for longer durations [9,37]. In few cases, Fe4N phase is also reported [13]. In the present investigation, there is no indication of CrN, suggesting that corrosion resistance of the steel is not impaired
Fig. 5. (a) TEM micrograph and (b) SAED pattern of severely shot peened sample.
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Fig. 6. (a) TEM micrograph and (b) corresponding SAED pattern from top surface of peened-nitrided sample.
as chromium has not diffused out from the matrix. This is distinctly different than investigations [9,37] in which extensive formation of CrN during nitriding is reported. Yu et al. [25] proposed a two- step transformation mechanism from martensite to expanded austenite. It suggests a complete transformation of strain induced martensite to austenite initially, followed by the transformation of austenite to S-phase. Several studies [12,15,24] have justified their results based on this theory. Further, a study by Ferkel et al. [17] indicates the presence of martensite in nitride layer. It suggests that the martensite could be due to the strain induced transformation of austenite during prior severe deformation or due to the decomposition of expanded austenite to martensite and CrN. The study by Ji et al. [14] confirmed the formation of strain induced martensite upon deformation, while phases formed after plasma nitriding were left unidentified suggesting that it could be a mixture of ferritic and/or martensitic phases along with Fe and/or Cr nitrides. This ambiguity was due to observed broad XRD peaks in the region of 40–45° and hence, TEM analysis is required to confirm the phases present. In a recent study [23] also, it has been stated that the presence of strain induced martensite enhances the nitriding layer thickness and it transforms to expanded austenite when nitrided at lower temperatures while it promotes CrN formation at higher treatment temperatures. In the present study it is confirmed that, under the process parameters
used for the duplex treatment, precipitation of deleterious CrN is not observed. It is also evident from the TEM studies that the major phase in the surface and subsurface layer is martensite. Hence, it is reasonable to assume that the nitrogen has entered into martensite lattice during nitriding treatment to facilitate the formation of about 45 μm thick nitride layer. However, no peak shift or broadening is observed in the XRD peaks of martensite due to nitrogen incorporation. Nitrogen being an austenite stabilizer, its supersaturation at the surface facilitates the transformation of martensite to expanded austenite in few regions as reported by Yu et al. [25]. It is evident by appearance of broad peak of S-phase in Fig.4 and weak diffraction rings of FCC-structure in Fig. 6(b). This study differs from most of the previous reports indicating that the major phase in the nitride layer is martensitic in nature. 5. Conclusion The study shows primarily that a nitride case depth of about 45 μm is successfully produced by plasma nitriding of 316L stainless steel at 400 ° C with shot peening pre-treatment. Nitriding case depth obtained is much higher compared to the reported literature for nitriding duration of 4 h. Shot-peening as a pre-treatment step induced nanocrystalline surface layers on the sample having martensite formed due to strain
Fig. 7. (a) TEM micrograph and (b) corresponding SAED pattern of peened-nitrided sample from about 15 μm below the surface.
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Fig. 8. Schematic diagram of formation of martensite during shot peening.
induced transformation of austenite. Nanocrystallinity is retained even after plasma nitriding treatment. Top layer of nitride is found to consist of majority of martensitic phase; while sub-surface (about 15 μm below the surface) found to possess lath martensite structure. Weak reflections of S-phase in XRD and diffraction rings corresponding to FCC structure is attributed to the transformation of a fraction martensite to Sphase, due to supersaturation of nitrogen in the martensite lattice. Hence it is inferred that the nitride layer is formed largely by incorporation of nitrogen into martensite lattice; contrary to majority of the reports. Also, CrN is completely absent due to the lower nitriding temperature employed in the study. Yet, thick nitride layer is obtained, which is attributed to extensive formation of defects and martensite during shot-peening pre-treatment. Acknowledgements We would like to acknowledge the Director, NITK and MHRD, Govt. of India for providing financial assistance for this research work. References [1] Y. Sun, T. Bell, Sliding wear characteristics of low temperature plasma nitrided 316 austenitic stainless steel, Wear 218 (1998) 34–42, http://dx.doi.org/10.1016/ S0043-1648(98)00199-9. [2] E. Menthe, A. Bulak, J. Olfe, A. Zimmermann, K.T. Rie, Improvement of the mechanical properties of austenitic stainless steel after plasma nitriding, Surf. Coat. Technol. 133–134 (2000) 259–263, http://dx.doi.org/10.1016/S0257-8972(00)00930-0. [3] E. Menthe, K.T. Rie, Further investigation of the structure and properties of austenitic stainless steel after plasma nitriding, Surf. Coat. Technol. 116-119 (1999) 199–204, http://dx.doi.org/10.1016/S0257-8972(99)00085-7. [4] W. Liang, Surface modification of AISI 304 austenitic stainless steel by plasma nitriding, Appl. Surf. Sci. 211 (2003) 308–314, http://dx.doi.org/10.1016/S01694332(03)00260-5. [5] Z.L. Zhang, T. Bell, Structure and corrosion resistance of plasma nitrided stainless steel, Surf. Eng. 1 (1985) 131–136, http://dx.doi.org/10.1179/sur.1985.1.2.131. [6] C. Blawert, A. Weisheit, B.L. Mordike, R.M. Knoop, Plasma immersion ion implantation of stainless steel: austenitic stainless steel in comparison to austenitic-ferritic stainless steel, Surf. Coat. Technol. 85 (1996) 15–27, http://dx.doi.org/10.1016/ 0257-8972(96)02880-0. [7] L. Gil, S. Brühl, L. Jiménez, O. Leon, R. Guevara, M.H. Staia, Corrosion performance of the plasma nitrided 316L stainless steel, Surf. Coat. Technol. 201 (2006) 4424–4429, http://dx.doi.org/10.1016/j.surfcoat.2006.08.081. [8] W.P. Tong, Nitriding iron at lower temperatures, Science 299 (2003) 686–688, http://dx.doi.org/10.1126/science.1080216 (80-.). [9] Y. Lin, J. Lu, L. Wang, T. Xu, Q. Xue, Surface nanocrystallization by surface mechanical attrition treatment and its effect on structure and properties of plasma nitrided AISI 321 stainless steel, Acta Mater. 54 (2006) 5599–5605, http://dx.doi.org/10.1016/j. actamat.2006.08.014. [10] M. Laleh, F. Kargar, M. Velashjerdi, Low-temperature nitriding of nanocrystalline stainless steel and its effect on improving wear and corrosion resistance, J. Mater. Eng. Perform. 22 (2013) 1304–1310, http://dx.doi.org/10.1007/s11665-012-0417-7. [11] M. Chemkhi, D. Retraint, A. Roos, C. Garnier, L. Waltz, C. Demangel, et al., The effect of surface mechanical attrition treatment on low temperature plasma nitriding of an austenitic stainless steel, Surf. Coat. Technol. 221 (2013) 191–195, http://dx.doi.org/ 10.1016/j.surfcoat.2013.01.047. [12] T. Balusamy, T.S.N.S. Narayanan, K. Ravichandran, I.S. Park, M.H. Lee, Plasma nitriding of AISI 304 stainless steel: role of surface mechanical attrition treatment, Mater. Charact. 85 (2013) 38–47, http://dx.doi.org/10.1016/j.matchar.2013.08.009.
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