In vitro corrosion fatigue behavior of low nickel high nitrogen austenitic stainless steel

In vitro corrosion fatigue behavior of low nickel high nitrogen austenitic stainless steel

Materials Science and Engineering A 538 (2012) 224–230 Contents lists available at SciVerse ScienceDirect Materials Science and Engineering A journa...

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Materials Science and Engineering A 538 (2012) 224–230

Contents lists available at SciVerse ScienceDirect

Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea

In vitro corrosion fatigue behavior of low nickel high nitrogen austenitic stainless steel Sudhakar Rao G a , Vakil Singh a,∗ , L.K. Singhal b a b

Department of Metallurgical Engineering, Institute of Technology, Banaras Hindu University, Varanasi 221005, U.P., India Director, R&D Division, Jindal Stainless Limited, Hisar, India

a r t i c l e

i n f o

Article history: Received 7 July 2011 Received in revised form 26 December 2011 Accepted 2 January 2012 Available online 18 January 2012 Keywords: In vitro corrosion fatigue LCF Low nickel stainless steel High nitrogen austenitic stainless steel Biomaterial Deformation induced martensite

a b s t r a c t Corrosion fatigue behavior of a new austenitic stainless steel, with relatively low nickel, high nitrogen and manganese (216L), and the usual 316L steel was studied in SBF, both in stress and strain control mode. Corrosion fatigue life of 216L in stress control mode was found to be longer to that of the 316L at ±450 MPa and there was opposite trend at high stress amplitudes. In strain control fatigue life of 216L was found longer than that of 316L both in air and in SBF over the strain amplitudes studied. The higher fatigue resistance of the 216L is attributed to high nitrogen content and higher stability of the austenite. © 2012 Elsevier B.V. All rights reserved.

1. Introduction Some metallic materials are considered highly promising for orthopaedic implants because of unique combination of their mechanical properties (strength, toughness, fatigue strength) and corrosion resistance, as compared to other biomaterials like polymers and ceramics. Metallic biomaterials are used in several medical devices such as artificial joints, bone plates, screws, intramedullary nails, stents, spinal fixations, spinal spacers, dental implants, wires, pace maker cases and artificial heart valves [1,2]. Stainless steels, cobalt-based alloys, commercially pure titanium (CP-Ti) and titanium alloys are known to be most widely used metallic biomaterials. However, there is growing interest in Ti and its alloys, particularly as dental and orthopaedic implants because of their relatively lower modulus of elasticity, superior biocompatibility and higher corrosion resistance than those of stainless steels and cobalt base alloys [3]. The 316L grade of stainless steel, however, is still the widely used material for internal fixation devices like fracture plates, screws, hip nail and THRs stems because of their good mechanical properties, acceptable biocompatibility, adequate corrosion resistance and low cost as compared with those of other metallic implant materials [4,5]. The major problem with use of

∗ Corresponding author. E-mail addresses: [email protected], [email protected] (V. Singh). 0921-5093/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2012.01.034

316L steel as bioimplant is from its high nickel content, which is known to have adverse effect on biocompatibility and several health related issues when used for longer durations. It has been established by several studies that metallic ions of nickel, chromium and cobalt, released during the process of corrosion in human body, tend to accumulate in many organs such as lungs, liver, kidneys and spleen [6,7]. Especially nickel ions are known to aggravate existing dermatitis and cause sarcoma, allergy [8–10]. Based on the studies on carcinogenic risks to human body from surgical implants and other foreign bodies by the International Agency for Research on Cancer (IARC) of the World Health Organization (WHO), it was estimated that nickel and its alloys are possibly carcinogenic to human body [10]. This element is also known for allergic sensitization especially in females. It is estimated that 30% of women suffer from skin allergy from the objects containing nickel [11,12]. All these problems caused by nickel and nickel containing alloys have led to development of low nickel or even nickel free austenitic stainless steels, replacing nickel by manganese and nitrogen. In recent years several investigations have been carried out on evaluation of mechanical properties, corrosion, corrosion fatigue and fatigue behavior of low nickel and nickel free high nitrogen austenitic stainless steels [13–26]. It is reported that addition of elemental nitrogen to austenitic stainless steels is beneficial in enhancing its resistance against pitting corrosion, improving fatigue resistance both in air as well as in chloride containing

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Table 1 Chemical composition of 216L and 316L stainless steels (wt%). Steel

C

Ni

Cr

Mn

N

Cu

Mo

Si

Ti

S

P

Fe

216L 316L

0.025 0.02

6.56 10.3

16.92 16.9

5.02 1.13

0.193 0.0358

1.94 0.303

1.75 2.02

0.23 0.483

– 0.007

0.005 0.003

0.032 0.0362

Bal. Bal.

environment. Further, the overall hazardous effect of nickel on human body should reduce with lowering of nickel content, using nitrogen as its substitute. The harmful effect of nickel in the austenitic stainless steel is eliminated when it is completely replaced by nitrogen and/or nitrogen/manganese. However, as the nitrogen level increases in these steels, their ductile to brittle transition temperature (DBTT) or brittleness also increases [27]. In view of this it is suggested to keep the level of nitrogen below 0.9% for medical applications to avoid ductile to brittle transition of the implant at body temperature [28]. Recently a new austenitic stainless steel (216L) has been developed by the M/s Jindal Steel Ltd. (JSL), Hisar, India, with relatively lower nickel content and higher nitrogen and manganese content than that of 316L steel, for medical application as a substitute of 316L. Since quite a few orthopaedic implants, particularly hip and leg are subjected to repeated loading for large number of cycles and usually fail due to corrosion fatigue, the present investigation was undertaken to characterize corrosion fatigue behavior of this newly developed 216L stainless steel, in simulated body fluid (SBF), and compare with that of the conventional 316L stainless steel. 2. Materials and methods The material of the present investigation, the newly developed austenitic stainless steel (216L), was obtained from the M/s JSL, Hisar, India, in the form of forged rods of 18 mm diameter and 250 mm length. Bars of 216L steel, forged at 1050 ◦ C, were annealed at 1050 ◦ C for 30 min for homogenization. Also the 316L stainless steel was supplied by the M/s. JSL in the form of plate of 203.2 × 203.2 × 16 mm size, annealed at 1050 ◦ C for 30 min. Chemical compositions of both the steels are recorded in Table 1. Optical microstructure of both the steels was characterized following mechanical polishing and etching with a solution of 15 ml HCl and 25 ml HNO3 . 2.1. Tensile testing Tensile properties of the two steels were determined using a 50 kN screw-driven InstronTM Universal Testing Machine. All the tensile test specimens were prepared as per the BS specification No. 12-1950, with gage length and diameter of 15.4 mm and 4.5 mm respectively. Tensile tests were conducted at a strain rat of 5 × 10−4 s−1 at room temperature. 2.2. Fatigue tests Axial push–pull stress controlled as well as strain controlled corrosion fatigue tests were carried out under fully reversed loading (R = −1) using 50 kN Servo hydraulic MTSTM (Model 810) under aerated simulated body fluid (SBF) environment. The composition of the SBF used is given in Table 2. Stress controlled corrosion fatigue tests were performed at different total stress amplitudes of ±400, ±450, ±500, ±550 and ±600 MPa, at a cyclic frequency of 1 Hz in

Fig. 1. Optical micrographs showing microstructures of the two steels in annealed condition (a) 216L (b) 316L.

SBF environment around gage section of the test specimen using a Perspex container attached to upper portion of the lower threaded end of the fatigue test specimen. Strain controlled low cycle corrosion fatigue tests were conducted in aerated SBF environment, at different total strain amplitudes (εt /2) ±0.50%, ±0.58%, ±0.64%, ±0.67% and ±0.73% at a constant strain rate of 5 × 10−3 s−1 . In order to bring out the effect of corrosive SBF environment on fatigue life, LCF tests were conducted also in air at respective total strain amplitudes at the same constant strain rate of 5 × 10−3 s−1 . In the strain controlled corrosion fatigue tests the SBF was made to drop continuously and slowly from circumference of upper shoulder of the specimen on cotton, wrapped around gage section of the LCF specimen. The lower threaded portion of the specimen was screwed in the lower grip, through a hole drilled in bottom of a disc shaped Perspex container, placed below the lower shoulder of the specimen, to collect the run-off solution. Fresh SBF was used for each test. All

Table 2 Composition of simulated body fluid (SBF) (g/l). NaCl

KCl

CaCl2

NaHCO3

Na2 HPO4 ·2H2 O

MgCl2 · 6H2 O

KH2 PO4

MgSO4 ·7H2 O

Glucose

8.00

0.40

0.18

0.35

0.48

0.10

0.06

0.10

1.00

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Fig. 2. S–N curve for the 216L and 316L stainless steels in simulated body fluid environment.

the tests were run until complete fracture. Fracture surfaces were examined using FESEM Quanta 200FE, following ultrasonic cleaning in acetone. X-ray diffraction of gage section adjacent to fracture surface of the fatigue tested samples was carried out to examine strain induced phase transformation, if any. TEM thin foils were also prepared from transverse section gage length of the fatigue tested sample, adjacent to the fracture surface, using twinjet electro polishing method in an electrolyte of 20% perchloric acid in methanol at 20 V and temperature of −30 ◦ C. 3. Results 3.1. Microstructure Optical microstructure of the 216L and 316L steels, in annealed condition, is shown in Fig. 1(a) and (b), respectively. As expected there is single phase in both the steels, with equiaxed, average grains size of 104 ␮m and 136 ␮m in the 216L and 316L steel respectively. Further, the distribution of grain size is quite uniform. 3.2. Tensile properties Tensile properties of the 216L and 316L steels in the annealed condition are presented in Table 3. It may be seen that yield strength of the 216L steel is higher by ∼30% than that of the 316L steel,

Fig. 4. Cyclic stress response of 216L steel in (a) air, (b) SBF at different total strain amplitudes.

however, tensile strength of the 216L steel is higher only marginally (∼5%). It is obvious that the rate as well as the degree of work hardening (UTS/YS) of the 316L steel is higher than that of the 216L. On the other hand there is opposite trend in ductility of the two steels. 3.3. Fatigue 3.3.1. Stress controlled fatigue The variation of fatigue life of the two steels with applied stress, in the aerated SBF environment is shown by the S–N plot in Fig. 2. It is evident from this plot that fatigue life of the 216L steel at low stress level (<±500 MPa) is nearly double of the 216L steel. However, there is opposite trend in their fatigue lives at higher stress amplitudes and a crossover occurs at stress level of approximately ±485 MPa.

Fig. 3. Coffin–Manson plots of the 216L and 316L steels in air and SBF.

3.3.2. Strain controlled fatigue LCF life of the two steels at different total strain amplitudes is brought out by Coffin–Manson (εp /2 vs 2 Nf ) plots on logarithmic scale (Fig. 3). It may be seen that fatigue life of the 216L steel is considerably longer than that of the 316L, over the entire range of the strain amplitude investigated. Fatigue life of the both the steels, may be seen to be considerably reduced under the SBF

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Table 3 Tensile properties of the 216L and 316L steels in annealed condition. Material

YS (MPa)

UTS (MPa)

Uniform plastic elongation (%)

Elongation (%)

True strain to fracture (␧f )

216L 316L

299 225

663 629

50 73

57 89

1.24 2.12

YS: Yield strength, UTS: Ultimate tensile strength.

environment. Further, the decrement in the fatigue life due to SBF environment increases with decrease in the strain amplitude. While the reduction in fatigue life of the 216L in SBF environment at the lowest total strain amplitude of ±0.50% is only 6%, it is more than 25% in the case of the 316L steel. The values of the different LCF parameters for both the steels in air as well as in SBF environment are recorded in Table 4. While the fatigue ductility exponent (c) for the 316L steel is close to the usual value of −0.5 to −0.7, it is much lower for the 216L steel. Also there is much difference in the values of the fatigue ductility coefficient (ε f ) of the two steels. It may be due to difference in deformation behavior of the two steels and limited fatigue test data over a narrow strain range, in particular in the 216L steel. Cyclic stress response of the 216L and 316L steels at different strain amplitudes in air and SBF environment are shown in

Fig. 6. X-ray diffraction profiles of fatigue sample tested (a) at total strain amplitude of ±0.64% for 216L and 316L and (b) at total stress amplitudes of ±400 MPa and ±500MPa for 316L steel.

Fig. 4(a) and (b) and Fig. 5(a) and (b), respectively. It may be seen that in general there is cyclic softening in the 216L steel followed by stabilization in the later stage except at the high strain amplitude of ±0.67% at which there is mild cyclic hardening in the beginning up to about 20 cycles. In contrast there is cyclic hardening from the very beginning up to about 20 cycles, followed by mild softening, stabilization, and pronounced hardening after ∼200 cycles till fracture in the 316L steel. The degree of cyclic hardening in the later stage increases whereas the extent of stabilization decreases, with increase in strain amplitude. Table 4 Low cycle fatigue parameters of 216L and 316L stainless steels in annealed condition, in air and SBF. Environment

216L 

Fig. 5. Cyclic stress response of 316L steel in (a) air, (b) SBF at different total strain amplitudes.

Air SBF

316L

εf

c

ε f

c

0.02104 0.03304

−0.1752 −0.2314

0.1209 0.7639

−0.41762 −0.65611

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Fig. 7. Optical micrographs showing slip pattern on cylindrical surface of the LCF specimens, close to fracture end, tested in SBF environment. 216L steel (a) ␧t /2 ±0.5% (b) ␧t /2 ±0.64% and 316L steel (c) εt /2 ±0.5% (d) ␧t /2 ±0.64%.

Fig. 8. TEM micrograph of the 216L steel tested in LCF at total strain amplitude (εt /2) ±0.64 (a) showing planarity of slip (b) corresponding diffraction pattern and (c) showing formation dislocation cell structure (d) corresponding diffraction pattern.

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3.3.3. Phase instability X-ray diffraction profiles of the LCF samples tested at εt /2 ± 0.64% for both the steels are shown in Fig. 6(a). It may be seen that while there are peaks of only austenite in the 216L, there are also peaks of ␣ -martensite along with austenite peaks in the 316L. Martensite peaks are due to strain induced phase transformation of metastable austenite to ␣ -martensite. X-ray diffraction profile of the 316L steel tested at different stress amplitudes, under stress controlled fatigue in the SBF environment are shown Fig. 6(b). It is evident that there are no martensite peaks in XRD profile of the 316L steel specimen tested at the lowest stress amplitude of ±450 MPa, whereas there are distinct ␣ -martensite peaks from the specimen tested at higher stress amplitude of ±500 MPa. Based on the fact that the intensity of the diffraction pattern of a particular phase in a mixture of phases depends on the concentration of that phase in the mixture [29], the volume fraction of martensite was calculated using the relationships (1) and (2). I␥ R␥ C␥ = I␣ R␣ C␣

(1)

C␥ + C␣ = 1

(2)

where I␥ and C␥ are integrated intensity and volume fraction of austenite phase respectively. I␣ and C␣ are integrated intensity and volume fraction of martensite phase respectively. R␥ and R␣ are constants and depend on the  and h k l values and the kind of substance (here it is for austenite and martensite). The volume fraction of the ␣ -martensite in the 316L sample tested at the stress amplitude of ±500 MPa was evaluated to be 10%. Volume fraction of ␣ -martensitein the specimen, tested at the lowest stress amplitude of ±400 MPa could not be estimated due to very small intensity of the martensite peak. It may be due to highly localized plastic deformation resulting from planar slip in a few slip bands at this stress level and consequent low volume fraction of strain induced ␣ -martensite in those slip bands. This observation is similar to earlier observation made by Nakajima et al. [30] and Topic et al. [31]. 3.3.4. Deformation behavior The morphology of the slip bands developed on cylindrical surface in gage section of the LCF specimen close to fracture end, tested in SBF environment was examined under optical microscope. The micrographs in Fig. 7(a) and (b) show slip pattern on 216L steel at total strain amplitudes of ±0.50 and ±0.64%, respectively. The surface features of the 316L steel at corresponding strain amplitudes are shown in Fig. 7(c) and (d), respectively. The planar nature of slip may clearly be seen in these micrographs. The TEM micrographs of the 216L steel, tested in strain controlled mode at εt /2 ± 0.64% are shown in Fig. 8 (a)–(d). Array of dislocations in planar slip band may clearly be seen in Fig. 8(a). Intersection of slip bands in other region resulting from multiple slip and formation of cell like structure may be seen in Fig. 8(c). 3.3.5. Fracture behavior Fracture behavior of the 216L and 316L steels resulting from strain controlled fatigue is shown in Figs. 9 and 10, respectively. It is obvious that there is marked difference in fatigue crack propagation behavior of the two steels as revealed by inter striation spacings in the fractographs. The inter striation spacing may be seen to be relatively less in the 216L than that in the 316L in respective test environments. The fatigue crack propagation rate was calculated for the 216L and 316L was found 0.869 ␮m/cycle and 1.78 ␮m/cycle in air and 1.54 ␮m/cycle and 2.40 ␮m/cycle in the SBF environment respectively. It may be noted that the rate of fatigue crack propagation in the 216L steel even at higher strain amplitude of ±0.67%

Fig. 9. SEM fractographs of the 216L steel, following LCF testing at total strain amplitude (εt /2) ±0.67%, in (a) Air (b) SBF.

is slower than that of the 316L steel even at lower strain amplitude of ±0.50%. 4. Discussion It is obvious from the tensile test data in Table 3 that yield strength of the 216L steel is longer by >30% than that of the 316L steel. However, the difference in their tensile strength is drastically reduced to ∼5% due to rapid and high degree of work hardening (UTS/YS) in the 316L steel. It may be attributed to combined effect of work hardening resulting from interactions of dislocations and increase in volume fraction of strain induced ␣ -martensite with strain. The higher yield strength of the 216L steel is essentially due to its high nitrogen content. Several reasons like solid solution strengthening, formation of interstitial solute complexes, clustering and ordering have been suggested for the resulting strengthening effect of nitrogen [30]. However, there is opposite behavior in ductility of the two steels. The larger fatigue life of the 216L than that of the 316L in stress control mode, at the low stress amplitude of 450 MPa, is due to higher strength of the former one. Relatively larger fatigue life of the 316L steel at high stress amplitude (≥ ±500 MPa) may be attributed to coaxing effect arising from rapid work hardening and the strain induced ␣ -martensite transformation [30,32,33]. The observations made in the present investigation on longer fatigue life of high nitrogen steel (216L) are in agreement with earlier observations made in high nitrogen containing stainless steels [22,23]. The detrimental role of corrosion arising from martensitic transformation in the 316L steel does not become effective at high stress amplitudes

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formed due to high nitrogen [25]. Nitrogen has also been reported to prevent dislocations from cross slip in the plastic zone around crack tips and retard fatigue crack growth [26]. The longer fatigue life of the 216L steel than that of the 316L steel may also be seen from the continuous cyclic softening in the former one. It is known that such cyclic softening suppresses the cyclic accumulation of fatigue damage resulting from the dislocation rearrangement from planar to cellular form [25]. A typical cellular structure may clearly be seen from the TEM micrograph in the 216L steel (Fig. 8 (c)). 5. Conclusions The following conclusions are drawn from the present investigation: 1. Fatigue life of the 216L stainless steel, in stress control mode in SBF environment is considerably longer than that of the 316L steel at stress amplitudes below ±500 MPa. However, at higher stress amplitudes the 316L steel exhibits longer fatigue life than that of the 216L. 2. Fatigue life of the 216Lsteel under total strain control mode is significantly longer than that of the 316L steel, both in air as well as in the SBF environment, over the entire range of the total strain amplitude investigated. 3. There is marked difference in cyclic stress response of the two steels both in air as well as in the SBF environment. While the 216L steel exhibits cyclic softening from beginning till the fracture, the 316L steel displays pronounced cyclic hardening in the later stage of cycling till fracture. References

Fig. 10. SEM fractographs of 316L steel, following LCF testing at total strain amplitude (εt /2) ±0.5%, in (a) air (b) SBF.

due to short exposures (0.36–1.47 h) of the specimens in the SBF environment during fatigue testing. Fatigue life of the 216L is significantly longer than that of the 316L over the entire range of the strain amplitude from ±0.50% to ±0.67%, both in air as well as in the SBF environment. It may be noted that disparity in fatigue lives of the two steels increases with decrease in strain amplitude. It is known that the major fraction of fatigue life in the high cycle fatigue is spent in the process of crack initiation, whereas in low cycle fatigue major fraction of life is spent in crack propagation. At low strain amplitude the applied plastic strain is accommodated in a few planar slip bands in favorably oriented grains. The plastic deformation in such bands causes ␣ -martensitic transformation in the 316L steel (Fig. 6). Thus the process of fatigue crack initiation becomes easier in the metastable 316L steel and consequently results in shorter fatigue life. The higher stability of the 216L steel against transformation of the austenite to ␣ -martensite offers higher resistance to crack initiation and leads to larger fatigue life. The shortening of fatigue life in the SBF environment in both the steels, in general, may be due to corrosive effect and hence early crack initiation. It may however be noted that nitrogen as solid solution in the austenitic steel increases its corrosion resistance. The effect of corrosive environment would be relatively more in the 316L steel because of ␣ -martensitic transformation in the slip bands and also in the plastic zones at the tip of the propagating cracks. Higher fatigue resistance of the nitrogen containing steels has been attributed to strong interaction between dislocations and nitrogen atoms as well as short range ordered zone

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