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Improvement of the blasting induced effects on medical 316 LVM stainless steel by short-term thermal treatments
Surface & Coatings Technology 258 (2014) 1075–1081
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Improvement of the blasting induced effects on medical 316 LVM stainless steel by short-term thermal treatments S. Barriuso a, M. Jaafar b, J. Chao a, A. Asenjo b, J.L. Gonzalez-Carrasco a,c,⁎ a b c
Centro Nacional de Investigaciones Metalúrgicas (CENIM-CSIC), Madrid, Spain Instituto de Ciencia de Materiales de Madrid (ICMM-CSIC), Cantoblanco 28049, Madrid, Spain Centro de Investigación Biomédica en Red en Bioingeniería, Biomateriales y Nanomedicina (CIBER-BBN), Madrid, Spain
a r t i c l e
i n f o
Article history: Received 20 April 2014 Accepted in revised form 8 July 2014 Available online 15 July 2014 Keywords: 316 LVM stainless steel Grit blasting Microstructure Magnetic properties Fatigue strength
a b s t r a c t Blasting of 316 LVM steel was performed by using alumina and zirconia particles that yield rough (up to Ra ~8 μm) and nearly smooth (up to Ra ~1 μm) surfaces, respectively. Besides roughening, the severe plastic deformation imposed during blasting yields subtle sub-surface microstructural changes such as grain size refinement, α′martensite formation, and work hardening. In this work we investigate the thermal treatment at 700 °C of the blasted steel, aimed to reverse the strain induced α′-martensite to austenite, with special emphasis in the correlation of the microstructural changes with the fatigue resistance. Blasting with rounded zirconia particles increases the fatigue resistance whereas the opposite effect is observed using alumina ones, which are more abrasive and cause a huge amount of embedded particles acting as stress raisers. It is shown that thermal treatment at 700 °C is very effective in the reversion of martensite, even after 2 min of annealing. Interestingly, the short thermal treatment of the alumina blasted steel increases the fatigue strength to values close to nonblasted material (400 MPa). Prolonged exposures, however, contribute to a full relaxation of the beneficial compressive residual stresses, being the net effect a decrease in the fatigue limit (340 MPa). © 2014 Elsevier B.V. All rights reserved.
1. Introduction Austenitic stainless steel 316 LVM (Low Vacuum Melting) is one of the most frequently used biomaterials for external or internal fixation devices because of a good combination of mechanical properties, biocompatibility and cost effectiveness. Moreover, the possibility of bending and shaping the implant makes it a desirable class of implant material to create a custom fit in the operating room [1]. Consequently, many times it is the material of choice for temporary osteosynthesis applications. Sometimes, however, it is considered as an attractive material, among others, for the fabrication of intramedullary nails for the proximal femur and diaphysary fractures or instrumented spinal arthrodesis to correct scoliosis, providing an optimal combination between high resistance during the consolidation period and a minimal invasive geometry. In these cases, the instrumentation is routinely maintained in situ for the patient lifetime. Unfortunately, a significant increase of metal ion release in the serum was demonstrated in patients with failed plates [2,3]. Recent studies consider that improved stabilisation of short term implants used for bone fracture repair may be crucial at the early-stage of implantation.
⁎ Corresponding author at: Centro Nacional de Investigaciones Metalúrgicas (CENIMCSIC), Madrid, Spain. Tel.: +34 915538900; fax: +34 915347425. E-mail address: [email protected] (J.L. Gonzalez-Carrasco).
http://dx.doi.org/10.1016/j.surfcoat.2014.07.027 0257-8972/© 2014 Elsevier B.V. All rights reserved.
A cost-effective method to improve the fixation and mechanical stability has been the development of a rough surface by grit blasting with oxide particles (mostly SiO2, ZrO2, or Al2O3), being the final roughness a function of the processing parameters (pressure, distance, time,…) and blasting particles (nature, shape, size). For instance, the angular particles lead to more material removal than rounded ones [4] and the incident angle influences noticeably the depth of the residual stress generated [5]. Therefore, the appropriated procedure should be employed accordingly with the specific application. Besides roughening of the surface, blasting of 316 LVM steel is known to induce a near surface work-hardened nanoscale microstructure [6], which will serve to retard fatigue-crack initiation by limiting plastic deformation, and compressive residual stresses [7], which will serve to retard fatigue-crack propagation. Both features play a beneficial role in terms of fatigue resistance, especially when considering the blasting with zirconia particles [4]. Relevant for the intended applications is that grit blasting of 316 LVM induces the subsurface formation of a very small volume fraction of strain induced α′-martensite [6,7], which is ferromagnetic. Despite that its presence must not be considered a limiting factor when applying magnetic resonance imaging (MRI) for clinical diagnosis [6], it could play a detrimental role for the corrosion resistance and ion release [8, 9], because the structural inhomogeneities would increase the density of the localized states [10]. The role of the strain induced martensite on the fatigue strength is limited to the previous works that study
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martensite formation during monotonic [11] or cycling [12,13] loading rather than to pre-deformed materials, as occurs in the blasted affected zone. Overall a beneficial role of martensite is found because, whereas cracks initiated exclusively in martensite regions, its formation results in an increase of the specific volume (~3%) and gives rise to compressive stresses, leading to premature contact of the crack surfaces and hence to a retardation of the short crack growth [14]. Therefore, advantages and disadvantages of applying additional thermal treatments for reversion of the strain induced martensite in blasted steel are open for discussion. The transformation of α′-martensite into austenite in metastable steels (ex. 301, 302, 304, 304L, 316 and 316L) is usually achieved by varying the starting and final temperatures of the reversion from 433 °C to 445 °C and from 705 °C to 724 °C, depending on the amount of deformation and heating rate [15,16]. The applicability of these thermal treatments to grit blasted 316 LVM steel could be, however, limited by the simultaneous relaxation of the sub-surface compressive residual stresses developed during blasting. Since fatigue cracks are likely to initiate within the blasted affected zones, it is important to characterize the nature and thermal stability of these regions. Although the stability of the residual stresses in mechanically surface-treated materials has been well studied, e.g., [17,18], little information exists on the corresponding thermal stability of work-hardened near-surface nanostructures [19]. In this work thermal treatments of blasted steel will be performed at 700 °C for 2 min and 1 h, which in addition to the reversion of martensite could provoke a partial and a full relaxation of the residual stresses, respectively [20]. 2. Experimental part For this investigation we have used specimens removed from a hot rolled and quenched bars (ϕ 30 mm) of 316 LVM austenitic stainless steel, which chemical composition (wt.%) is Cr 17.48, Ni 14.13, Mo 2.87, Mn 1.62, Si 0.53, C 0.024, Cu 0.067, N 0.061, S 0.001, and Fe in balance. For microstructural characterization purposes, discs of about 20 mm in diameter and 2 mm thick were removed perpendicularly to the rolling direction of the bar and prepared by conventional metallographical techniques. For the rotating bending fatigue tests, a set of specimens of 80 mm in length with 5.89 mm diameter in the gauge length was machined from rods that were removed parallel to the rolling direction of the bar. Blasting of the discs and fatigue specimens was performed by the implant manufacturer (SURGIVAL SL, Valencia, Spain) with two different types of particles under a pressure of 350 KPa for 1 min and with a distance between the nozzle and the target surface of about 20 cm. A first set of samples, hereinafter BL_ZrO, has been blasted using zirconia rounded particles sized between 125 μm and 250 μm. The second set of samples, hereinafter BL_AlO, has been blasted with alumina angular particles (white corundum) sized between 1 mm and 2 mm. Whereas discs were fully blasted, fatigue specimens were only blasted at the gauge length. All specimens were cleaned following the standard procedures and passivated in acid citric before delivered by the implant manufacturer. Thermal treatments were performed by introducing the specimens into a furnace at the test temperature (700 °C) for 2 min and 1 h. After treatments, the specimens were removed from the furnace and left to cool down to ambient temperature. Quantitative surface roughness was determined with a profilometer Mitutoyo Surftest 401. The measurements were obtained from line profiles along a 4 mm length. The surface roughness was characterized by average surface roughness (Ra) in micrometer at a high sensitivity setting (0.01 μm). Microstructural characterization of the surface morphology and cross sections of the specimens was carried out by using a scanning electron microscope (JEOL-6500F) equipped with a field emission gun (FEG) and coupled with an energy dispersive X-ray (EDX) system for
chemical analysis. The microstructure was revealed by Backscattered Electron Images (BEI) obtained on fresh ground and polished surfaces at low voltages. Contrast of the image is only associated with the different crystallographic orientation of the grains since in this case heterogeneities in composition are not expected. Average grain size was determined on a set of representative fields of view at 2000 magnifications. Phase identification at the blasted affected zones was performed by the X-ray diffraction (XRD) technique. Measurements were carried out with a Bruker AXS D8 diffractometer equipped with Co tube and Goebel mirror optics to obtain a parallel and monochromatic X-ray beam. A current of 30 mA and a voltage of 40 kV in both, grazing incidence condition and conventional θ–2θ scan were used. The evolution of the α′-martensite during these thermal treatments will be studied by magnetic force microscopy (MFM) on polished cross sections, which allows the detection of this phase even in nanocrystalline size without the need of any special preparation. Combined atomic force microscope (AFM) and magnetic force microscope (MFM) imaging have been performed using a microscope from Nanotec Electrónica S.L. The topography of the surface (AFM) and the magnetic image are acquired simultaneously in dynamic mode [21]. Nanosensors standard and homemade low moment MFM probes with a force constant of 3 N/m and a resonance frequency of 75 kHz have been used in these experiments [22]. Zones of about 40 μm in depth from the blasted surfaces were scanned. The hardness distribution beneath the surface was measured on cross-sections of blasted specimens previously coated with a relatively thick layer of Cu obtained by electrolytic deposition. This coating preserves the blasted surface during preparation of the cross section but also allows the approach to the initial blasted interface with the indenter. To avoid interactions between adjacent indentations the measurements were performed perpendicularly to the blasted interphase drawing a zigzag. The Vickers microhardness measurements were performed in Wilson equipment using 15 s of dwell time. Load was very small (98 mN) to obtain the indentations of small size and then approach to the blasted surface as much as possible. Fatigue specimens were tested in a rotating bending fatigue machine by using a symmetrical ratio and alternative cycle of stresses (R = −1). The specimens were tested with constant stress amplitude at different load levels. The fatigue strength was established as the maximum stress after 107 cycles without failure. The fracture surfaces of tested samples were also characterized by SEM. 3. Results and discussion 3.1. Microstructural characterization Blasting of austenitic stainless steel with either alumina or zirconia particles produces a severe surface plastic deformation that yields an irregular rough surface morphology with an average Ra of about 8 and 1 μm, respectively. These differences should be understood when considering the larger size and angular shape of the alumina particles that severely contribute to grind down the material. This abrasive role is also manifested when considering the mass loss (0.0076 g) when blasting with alumina particles, whereas a mass increase (0.0047 g) occurs when blasting with zirconia. Heat treatment of the samples causes a change in the aspect of the surface from a glazed grey to a rather dark blue colour due to the moderated oxidation process. Mass changes in the un-blasted and blasted conditions after heat treatment, however, were not significant, which indicates the formation of a very thin oxide layer. As a matter of fact, surface roughness was also similar. SEM examination of the blasted and thermally treated surfaces, Fig. 1, reveals the presence of oxide particles of heterogeneous size, obviously remnants of the oxide blasting particles. X-Ray diffraction analysis indicates the presence of a moderated volume fraction of embedded alumina particles (~ 10%), many of them not visible by SEM. After
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a) Al2O3
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b)
ZrO
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relative large fragments of the original blast particles, which remain embedded in the matrix, are observed. Following the ultrafine-grained zone, highly deformed grains containing thin twins, likely also containing some strain induced α′-martensite, [7] are found (see insets number 2). The grains are not altered for depth of about 200 μm and 100 μm for alumina and zirconia blasted samples, respectively, where microstructure of the bulk is observed. Grain pattern of the bulk (bottom part) is characterized by a larger grain size (~ 25 μm) and the presence of wider thermally induced twins, obviously produced during the thermomechanical processing of the bar. Occasionally, elongated inclusions of MnS that are parallel to the longitudinal bar direction are observed. With regard to the effects of annealing on the microstructure of blasted specimens it is worth to remind that the recrystallization temperature is typically 1/3 to 1/2 of the melting temperature (can be as high as 0.7 Tm in some alloys). The recrystallization temperature, however, decreases as the cold work is increased, although below a “critical deformation” recrystallization does not occur. Taking into consideration these issues and the gradient in the plastic deformation during blasting, a non-homogeneous evolution of the microstructure during annealing is observed. After 2 min of annealing, Fig. 3, the nanograin size pattern is now more evident, which is consistent with an initial recovery step that allows a better crystallographic contrast of the image. After 1 h of annealing, signs of recrystallization are evident at the outermost part of the blasted affected zones, where the higher amount of cold work caused a larger decrease in the recrystallization temperature. As can be seen in Fig. 4, recrystallization causes the formation of new grains and a further growth achieving up to 1 μm in size, whereas the microstructure of the inner and less strained zone remains stable.
Fig. 1. Surface SEM images of specimens blasted with a) alumina and b) zirconia particles.
3.2. Magnetic force microscopy (MFM) thermal treatment, and especially after 1 h exposure, additional submicrometric particles, likely Fe- and Cr-rich oxides, which are not contributed to a roughness increase, are observed. Cross sectional examination, Fig. 2, reveals a grain size refinement beneath the blasted surface (see insets number 1), as expected [6,7]. Essentially, grain size ranges between 100 and 500 nm in a layer of about 30 μm for the BL-AlO samples, Fig. 2a, and 10–15 μm for the BL-ZrO ones, Fig. 2b. In the case of the specimens blasted with alumina, Fig. 2a,
XRD was used for phase identification but martensite peaks were not detected. MFM, which is more sensitive, revealed this phase and also allowed to study its distribution in the near surface region. A zone of about 40 μm in depth from the blasted surface was investigated. In the MFM images, dark and bright contrasts appear in certain regions of the samples. Dark contrast corresponds to attractive tip-sample interactions, while bright contrast can be associated with the regions where there is a repulsive tip-sample interaction.
Fig. 2. Cross sectional views of blasting affected zones corresponding to specimens blasted with a) alumina and b) zirconia particles. Insets (1) denote zones close to the blasted surfaces and inset (2) zones beneath the fine grained zone. SMn corresponds to inclusions of manganese sulphide.
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Fig. 3. Cross sectional views of blasting affected zones after thermal treatment at 700 °C for 2 min. a) Alumina blasted surfaces and b) zirconia blasted surfaces. Insets are close-up of zones beneath the blasted surfaces.
Fig. 4. Cross sectional views of blasting affected zones after thermal treatment at 700 °C for 1 H. a) Alumina blasted surfaces and b) zirconia blasted surfaces. Insets are close-up of zones beneath the blasted surfaces.
Fig. 5 shows typical topographic (Fig. 5a, b, c) and magnetic images (Fig. 5d, e, f) obtained on the cross sections, near to the surface, of the non-annealed (Fig. 5a, d) and annealed (Fig. 5b, c, e, f) zirconia blasted specimens. Comparison of Fig. 5a and d, corresponding to the same zone of the non-annealed and blasted specimen, indicates the lack of correlation between topography and magnetic signal, i.e. we can assume that the polishing process is not responsible for the magnetic behaviour. For the as-blasted condition, Fig. 5d, submicrometric α′martensite coexists with austenite characterized by the absence of magnetic signal. α′-Martensite is heterogeneously distributed along parallel and perpendicular directions to the surface. As we get closer to the surface, higher density of the magnetic phase is observed. Along the parallel direction, areas of high density of magnetic phase intercalate with areas of a lower density. In the transition area from the ultrafine to the non-affected microstructure typical α′-martensite needles can be found. Such lines present bright and dark contrast that corresponds to multidomain regions. In addition, dark spots that correspond to single domain crystals can also be found. Some “magnetic zones” present single domain configuration with low coercive fields, i.e. the magnetization in these regions will be always oriented parallel to the tip stray field, presenting homogeneous dark contrast. Thermal treatment of the BL_ZrO samples considerably reduces the amount of α′martensite and after 2 min, Fig. 5e, only some magnetic signal is observed at 15–20 μm from the surface. After 1 h of exposure, Fig. 5f, no magnetic signal is detected. In the BL_AlO sample, the MFM analysis reveals a more heterogeneous distribution of martensite along the blasted surface, with magnetic contrast at deeper regions, even at ~20 μm from the surface (images not shown). It is worth noting that the cumulative impact effect is necessary to produce α′-martensite formation, since a given threshold
strain level is needed for the transformation. The largest Al2O3 particles are more energetic than the ZrO2 particles and, therefore, lead to more cumulative deformation, which implies the formation of α′-martensite at deeper regions from the surface. On the other hand, their angular shape made that, depending of the attack angle to the surface, impact energy of some of them is devoted to tearing off material, leads to a micro-cutting effect, rather than to accumulate deformation. Hence, some of the α′-martensite accumulated from the previous impacts (which do not erode material) is also eliminated from the surface. Obviously, this increasing transformation process is retarded when local material removal by the large and edge-like Al2O3 particles occurs. Evolution of martensite with increasing exposure time is similar to that described for the BL-ZrO samples. Thermal treatment is very effective in causing a reversion of the martensite, even after only 2 min of exposure. This situation is illustrated in Fig. 6 that shows very small zones with magnetic contrast, Fig. 6b, around a large embedded alumina particle. 3.3. Microhardness Microhardness measurements of the blasted specimens, Fig. 7, reveal a gradient in hardness perpendicular to the blasted surface with a maximum of ~ 360 HV0.01 close to the surface, irrespectively the particle used for blasting and the thermal condition. In all cases, hardness achieves a near constant value of about 180 HV0.01 at a depth of ~ 200 μm and ~ 100 μm into the bulk, depending on whether the blasting was performed with particles of alumina, Fig. 7a, or zirconia, Fig. 7b, respectively. Thermal treatments at 700 °C do not significantly influence the hardness gradients, although depending on the relative depth to the
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Fig. 5. Topographic (a, b, c) and magnetic (d, e, f) images obtained by AFM (a, b, c) and MFM (d, e, f) in the cross sections of BL_ZrO samples in the as-blasted condition (a, d) and after annealing at 700 °C for 2 min (b, e) and 1 h (c, f). The line mark at the right indicates the edge of the sample (blasted surface).
a)
b)
Al2O3
around 100 μm, for the zirconia blasted specimens, are irrelevant, which is consistent with the thermal stability of the bulk microstructure at this temperature. 3.4. Rotating bending fatigue tests In a previous work [4] it has been shown that fatigue resistance of the 316 LVM steel (~400 MPa) is very close to the yield strength value (430 MPa). Interestingly, whereas blasting with zirconia rounded
Microhardness HV 0.01
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Non-treated 700ºC / 2 min 700ºC / 1 h
300 250 200
400
10
100 30 Depth (µm)
BL_ZrO
200
Non-treated 700ºC / 2 min 700ºC / 1 h
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Fig. 6. Topographic (a) and magnetic (b) images obtained by AFM (a) and MFM (b) in the cross sections of BL_AlO samples after annealing at 700 °C for 1 h. The line mark at the right indicates the edge of the sample (blasted surface) and the circles delimit zones of magnetic signal.
BL_AlO
350
150
Microhardness HV 0.01
blasted surface slight changes in hardness are observed. Fig. 7 shows hardness values of blasted specimens with and without heat treatment at depths of about 10, 30, 100 and 200 μm from the blasted surface. The differences in hardness with regard to the bulk are likely related to the presence of α′-martensite, which has higher hardness than the austenite [23], the grain refinement, as the Hall–Petch expression predicts, and work hardening. As can be seen, thermal treatment for 2 min causes a moderated decrease in hardness. Despite the fact that both samples have approximately the same quantity of α′-martensite phase, the decrease in hardness for the alumina blasted specimens is more relevant at around 100 μm from the surface, which is consistent with the presence of less areas showing magnetic contrast close to the surface (Fig. 6) and a lesser dislocation density. Further increase in hardness with increasing exposure for a given depth can be associated with the grain size refinement associated with the reversion of the α′-martensite. According to Lee et al. [24] during isothermal holding, the equiaxed grains are nucleated and grown due to the diffusive reverse transformation, resulting in equiaxed grains, additional austenite volume fraction, and compressive transformation strain. It is worth to notice that grain growth observed at the outermost part of the blasted affected zones, Fig. 4a and b, is out the zone of measurements. Variations in hardness at around 200 μm in depth, for the alumina blasted specimens, and at
10
30 100 Depth (µm)
200
Fig. 7. Subsurface hardness as a function of depth from the surface for the specimens blasted with a) alumina and b) zirconia particles.
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particles causes a moderated increase in the fatigue limit (~420 MPa), blasting with angular alumina particles yields a significant decrease (~360 MPa). The higher roughness and the presence of a high amount of embedded alumina particles, acting as stress concentrators, make the alumina blasted surface prone to crack nucleation and growth despite the subsurface compressive residual stresses and grain size refinement. Thermal treatment of the blasted specimens at 700 °C for 2 min and 1 h influences the fatigue limit, as summarised in Table 1. In the case of the specimens blasted with zirconia particles, short term annealing hardly affects the fatigue strength. After 1 h of annealing, however, a moderated decrease in the fatigue strength (b5%) is observed, achieving values typical of the non-blasted condition. Interestingly, thermal treatment of the alumina blasted specimens causes an increase in the fatigue strength (~ 8%) after short annealing times (2 min), Fig. 8, achieving values for fatigue strength of about 400 MPa. The long annealing treatment, however, decreases the fatigue strength to minimum values (340 MPa). The beneficial effect of the short annealing treatment on the fatigue behaviour of the alumina blasted condition must be analysed by considering the microstructural induced changes influencing the crack initiation process [25]. Due to the very short annealing, the compressive residual stresses developed during blasting, which will serve to retard fatigue-crack propagation, are only partially released, as determined for similar samples by non-contacting thermoelectric measurements [20]. It is worth mentioning that partial relaxation of the compressive residual stresses would be counterbalanced by the reversion of α′-martensite and the induced grain size refinement of the microstructure [26], which would play a beneficial role. It could be argued that zones experiencing a plastic strain during fatigue, obviously at the outermost zone, could develop new strain induced martensite. Although cracks would develop around alumina particles acting as stress concentrators [24], the nanograin size of austenite is likely being controlling the damage process by acting as indirect barriers for crack propagation. The crack which meets the first grain boundary is shorter in the fine microstructure and would induce less strain at the crack tip, minimizing the formation of new martensite. A question still arises about why the beneficial effect of the short term annealing is not manifested for the zirconia blasted specimens. Blasted affected zone is larger for the alumina blasted specimens. Moreover, from the analysis of the gradient in hardness it follows that plastic deformation is more severe when blasting with alumina particles. Therefore, a larger amount of strain induced martensite is expected to form in the alumina samples. The detrimental role of the longest exposure can be associated with the full release of the initial compressive residual stresses, [20]. The lower fatigue strength of the thermally treated steel blasted with alumina particles can be related to their lower thermal expansion coefficient with regard to that of the matrix where they reside. During cooling from the annealing temperature, textural residual stresses become compressive in the ceramic particles but tensile in the matrix [27], favouring crack formation. From the above discussion it follows that short-term thermal treatments of grit blasted austenitic stainless steels may be a useful tool to reverse α′-martensite developed at the blasted affected zone as consequence of the severe plastic deformation, resulting in an increase in the fatigue strength. Improvement on the corrosion behaviour and ion Table 1 Rotating (R = −1) fatigue limit (MPa) of alumina and zirconia blasted specimens with and without post thermal treatments. Material condition
Non-treated
Polished BL-ZrO BL-AlO
400 420 370
Thermally treated (700 °C) 2 min
1h
– 420 390–400
– 390–400 340
600 550
Stress (MPa)
1080
500 450 400 350 300
SG-Alumina o SG-Alumina + 700 C/2min o SG-Alumina + 700 C/1h
250 200 4
5
10
10
6
10
7
10
8
10
Number of cycles to failure Fig. 8. Stress amplitude versus number of cycles to failure during rotating bending tests of alumina blasted steel before and after annealing for 2 min and 1 h.
release, which are dominated by the surface properties, cannot be discarded. Therefore, in vitro experiments considering short and medium periods of immersion in simulated human fluids are in progress. 4. Conclusions From the above discussion the following conclusions can be drawn: • Grit blasting of 316 LVM austenitic steel with either alumina or zirconia particles causes a severe surface plastic deformation that roughs the surface, decreases grain size, and favours the formation of martensite, which is in agreement with the previous works. • Annealing at 700 °C reverse the strain induced martensite (ferromagnetic) at the blasted affected zone to austenite (paramagnetic) without compromising the stability of the bulk microstructure, which eliminates the eventual problem associated with the use of strong magnetic fields during clinical diagnostics. • Blasting with rounded zirconia particles increases the fatigue resistance and short term annealing (2 min) hardly affects such increase. After 1 h of annealing, however, a moderated decrease in the fatigue strength (b5%) is observed, achieving values typical of the nonblasted condition. • Blasting with alumina particles, which are more abrasive and causes a huge amount of embedded particles, causes a significant decrease in the fatigue strength. Interestingly, short annealing times (2 min) cause an increase in the fatigue strength (~ 8%) as consequence of the reversion of the martensite, achieving values corresponding to the non-blasted condition. Longer periods of annealing cause a moderate grain size increase at the outermost surface and contribute to full relaxation of the compressive residual stresses, yielding a significant reduction in the fatigue strength. Conflict of interest None. Acknowledgements The authors are thankful for financial support from the MICINN (MAT2009-14695-C04) and the MINECO (MAT 2012-37736-C05-01). S. Barriuso thanks to JAE predoc grant of CSIC. Technicians of Laboratories of X-Ray Diffraction, Mechanical Testing, and Microscopy (CENIM) and the SPM-Service of the ICMM-CSIC are specially acknowledged. References [1] J.A. Disegi, L. Eschbach, Injury 31 (Suppl. 4) (2000) 2–6.
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Improvement of the blasting induced effects on medical 316 LVM stainless steel by short-term thermal treatments S. Barriuso a, M. Jaafar b, J. Chao a, A. Asenjo b, J.L. Gonzalez-Carrasco a,c,⁎ a b c
Centro Nacional de Investigaciones Metalúrgicas (CENIM-CSIC), Madrid, Spain Instituto de Ciencia de Materiales de Madrid (ICMM-CSIC), Cantoblanco 28049, Madrid, Spain Centro de Investigación Biomédica en Red en Bioingeniería, Biomateriales y Nanomedicina (CIBER-BBN), Madrid, Spain
a r t i c l e
i n f o
Article history: Received 20 April 2014 Accepted in revised form 8 July 2014 Available online 15 July 2014 Keywords: 316 LVM stainless steel Grit blasting Microstructure Magnetic properties Fatigue strength
a b s t r a c t Blasting of 316 LVM steel was performed by using alumina and zirconia particles that yield rough (up to Ra ~8 μm) and nearly smooth (up to Ra ~1 μm) surfaces, respectively. Besides roughening, the severe plastic deformation imposed during blasting yields subtle sub-surface microstructural changes such as grain size refinement, α′martensite formation, and work hardening. In this work we investigate the thermal treatment at 700 °C of the blasted steel, aimed to reverse the strain induced α′-martensite to austenite, with special emphasis in the correlation of the microstructural changes with the fatigue resistance. Blasting with rounded zirconia particles increases the fatigue resistance whereas the opposite effect is observed using alumina ones, which are more abrasive and cause a huge amount of embedded particles acting as stress raisers. It is shown that thermal treatment at 700 °C is very effective in the reversion of martensite, even after 2 min of annealing. Interestingly, the short thermal treatment of the alumina blasted steel increases the fatigue strength to values close to nonblasted material (400 MPa). Prolonged exposures, however, contribute to a full relaxation of the beneficial compressive residual stresses, being the net effect a decrease in the fatigue limit (340 MPa). © 2014 Elsevier B.V. All rights reserved.
1. Introduction Austenitic stainless steel 316 LVM (Low Vacuum Melting) is one of the most frequently used biomaterials for external or internal fixation devices because of a good combination of mechanical properties, biocompatibility and cost effectiveness. Moreover, the possibility of bending and shaping the implant makes it a desirable class of implant material to create a custom fit in the operating room [1]. Consequently, many times it is the material of choice for temporary osteosynthesis applications. Sometimes, however, it is considered as an attractive material, among others, for the fabrication of intramedullary nails for the proximal femur and diaphysary fractures or instrumented spinal arthrodesis to correct scoliosis, providing an optimal combination between high resistance during the consolidation period and a minimal invasive geometry. In these cases, the instrumentation is routinely maintained in situ for the patient lifetime. Unfortunately, a significant increase of metal ion release in the serum was demonstrated in patients with failed plates [2,3]. Recent studies consider that improved stabilisation of short term implants used for bone fracture repair may be crucial at the early-stage of implantation.
⁎ Corresponding author at: Centro Nacional de Investigaciones Metalúrgicas (CENIMCSIC), Madrid, Spain. Tel.: +34 915538900; fax: +34 915347425. E-mail address: [email protected] (J.L. Gonzalez-Carrasco).
http://dx.doi.org/10.1016/j.surfcoat.2014.07.027 0257-8972/© 2014 Elsevier B.V. All rights reserved.
A cost-effective method to improve the fixation and mechanical stability has been the development of a rough surface by grit blasting with oxide particles (mostly SiO2, ZrO2, or Al2O3), being the final roughness a function of the processing parameters (pressure, distance, time,…) and blasting particles (nature, shape, size). For instance, the angular particles lead to more material removal than rounded ones [4] and the incident angle influences noticeably the depth of the residual stress generated [5]. Therefore, the appropriated procedure should be employed accordingly with the specific application. Besides roughening of the surface, blasting of 316 LVM steel is known to induce a near surface work-hardened nanoscale microstructure [6], which will serve to retard fatigue-crack initiation by limiting plastic deformation, and compressive residual stresses [7], which will serve to retard fatigue-crack propagation. Both features play a beneficial role in terms of fatigue resistance, especially when considering the blasting with zirconia particles [4]. Relevant for the intended applications is that grit blasting of 316 LVM induces the subsurface formation of a very small volume fraction of strain induced α′-martensite [6,7], which is ferromagnetic. Despite that its presence must not be considered a limiting factor when applying magnetic resonance imaging (MRI) for clinical diagnosis [6], it could play a detrimental role for the corrosion resistance and ion release [8, 9], because the structural inhomogeneities would increase the density of the localized states [10]. The role of the strain induced martensite on the fatigue strength is limited to the previous works that study
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martensite formation during monotonic [11] or cycling [12,13] loading rather than to pre-deformed materials, as occurs in the blasted affected zone. Overall a beneficial role of martensite is found because, whereas cracks initiated exclusively in martensite regions, its formation results in an increase of the specific volume (~3%) and gives rise to compressive stresses, leading to premature contact of the crack surfaces and hence to a retardation of the short crack growth [14]. Therefore, advantages and disadvantages of applying additional thermal treatments for reversion of the strain induced martensite in blasted steel are open for discussion. The transformation of α′-martensite into austenite in metastable steels (ex. 301, 302, 304, 304L, 316 and 316L) is usually achieved by varying the starting and final temperatures of the reversion from 433 °C to 445 °C and from 705 °C to 724 °C, depending on the amount of deformation and heating rate [15,16]. The applicability of these thermal treatments to grit blasted 316 LVM steel could be, however, limited by the simultaneous relaxation of the sub-surface compressive residual stresses developed during blasting. Since fatigue cracks are likely to initiate within the blasted affected zones, it is important to characterize the nature and thermal stability of these regions. Although the stability of the residual stresses in mechanically surface-treated materials has been well studied, e.g., [17,18], little information exists on the corresponding thermal stability of work-hardened near-surface nanostructures [19]. In this work thermal treatments of blasted steel will be performed at 700 °C for 2 min and 1 h, which in addition to the reversion of martensite could provoke a partial and a full relaxation of the residual stresses, respectively [20]. 2. Experimental part For this investigation we have used specimens removed from a hot rolled and quenched bars (ϕ 30 mm) of 316 LVM austenitic stainless steel, which chemical composition (wt.%) is Cr 17.48, Ni 14.13, Mo 2.87, Mn 1.62, Si 0.53, C 0.024, Cu 0.067, N 0.061, S 0.001, and Fe in balance. For microstructural characterization purposes, discs of about 20 mm in diameter and 2 mm thick were removed perpendicularly to the rolling direction of the bar and prepared by conventional metallographical techniques. For the rotating bending fatigue tests, a set of specimens of 80 mm in length with 5.89 mm diameter in the gauge length was machined from rods that were removed parallel to the rolling direction of the bar. Blasting of the discs and fatigue specimens was performed by the implant manufacturer (SURGIVAL SL, Valencia, Spain) with two different types of particles under a pressure of 350 KPa for 1 min and with a distance between the nozzle and the target surface of about 20 cm. A first set of samples, hereinafter BL_ZrO, has been blasted using zirconia rounded particles sized between 125 μm and 250 μm. The second set of samples, hereinafter BL_AlO, has been blasted with alumina angular particles (white corundum) sized between 1 mm and 2 mm. Whereas discs were fully blasted, fatigue specimens were only blasted at the gauge length. All specimens were cleaned following the standard procedures and passivated in acid citric before delivered by the implant manufacturer. Thermal treatments were performed by introducing the specimens into a furnace at the test temperature (700 °C) for 2 min and 1 h. After treatments, the specimens were removed from the furnace and left to cool down to ambient temperature. Quantitative surface roughness was determined with a profilometer Mitutoyo Surftest 401. The measurements were obtained from line profiles along a 4 mm length. The surface roughness was characterized by average surface roughness (Ra) in micrometer at a high sensitivity setting (0.01 μm). Microstructural characterization of the surface morphology and cross sections of the specimens was carried out by using a scanning electron microscope (JEOL-6500F) equipped with a field emission gun (FEG) and coupled with an energy dispersive X-ray (EDX) system for
chemical analysis. The microstructure was revealed by Backscattered Electron Images (BEI) obtained on fresh ground and polished surfaces at low voltages. Contrast of the image is only associated with the different crystallographic orientation of the grains since in this case heterogeneities in composition are not expected. Average grain size was determined on a set of representative fields of view at 2000 magnifications. Phase identification at the blasted affected zones was performed by the X-ray diffraction (XRD) technique. Measurements were carried out with a Bruker AXS D8 diffractometer equipped with Co tube and Goebel mirror optics to obtain a parallel and monochromatic X-ray beam. A current of 30 mA and a voltage of 40 kV in both, grazing incidence condition and conventional θ–2θ scan were used. The evolution of the α′-martensite during these thermal treatments will be studied by magnetic force microscopy (MFM) on polished cross sections, which allows the detection of this phase even in nanocrystalline size without the need of any special preparation. Combined atomic force microscope (AFM) and magnetic force microscope (MFM) imaging have been performed using a microscope from Nanotec Electrónica S.L. The topography of the surface (AFM) and the magnetic image are acquired simultaneously in dynamic mode [21]. Nanosensors standard and homemade low moment MFM probes with a force constant of 3 N/m and a resonance frequency of 75 kHz have been used in these experiments [22]. Zones of about 40 μm in depth from the blasted surfaces were scanned. The hardness distribution beneath the surface was measured on cross-sections of blasted specimens previously coated with a relatively thick layer of Cu obtained by electrolytic deposition. This coating preserves the blasted surface during preparation of the cross section but also allows the approach to the initial blasted interface with the indenter. To avoid interactions between adjacent indentations the measurements were performed perpendicularly to the blasted interphase drawing a zigzag. The Vickers microhardness measurements were performed in Wilson equipment using 15 s of dwell time. Load was very small (98 mN) to obtain the indentations of small size and then approach to the blasted surface as much as possible. Fatigue specimens were tested in a rotating bending fatigue machine by using a symmetrical ratio and alternative cycle of stresses (R = −1). The specimens were tested with constant stress amplitude at different load levels. The fatigue strength was established as the maximum stress after 107 cycles without failure. The fracture surfaces of tested samples were also characterized by SEM. 3. Results and discussion 3.1. Microstructural characterization Blasting of austenitic stainless steel with either alumina or zirconia particles produces a severe surface plastic deformation that yields an irregular rough surface morphology with an average Ra of about 8 and 1 μm, respectively. These differences should be understood when considering the larger size and angular shape of the alumina particles that severely contribute to grind down the material. This abrasive role is also manifested when considering the mass loss (0.0076 g) when blasting with alumina particles, whereas a mass increase (0.0047 g) occurs when blasting with zirconia. Heat treatment of the samples causes a change in the aspect of the surface from a glazed grey to a rather dark blue colour due to the moderated oxidation process. Mass changes in the un-blasted and blasted conditions after heat treatment, however, were not significant, which indicates the formation of a very thin oxide layer. As a matter of fact, surface roughness was also similar. SEM examination of the blasted and thermally treated surfaces, Fig. 1, reveals the presence of oxide particles of heterogeneous size, obviously remnants of the oxide blasting particles. X-Ray diffraction analysis indicates the presence of a moderated volume fraction of embedded alumina particles (~ 10%), many of them not visible by SEM. After
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a) Al2O3
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relative large fragments of the original blast particles, which remain embedded in the matrix, are observed. Following the ultrafine-grained zone, highly deformed grains containing thin twins, likely also containing some strain induced α′-martensite, [7] are found (see insets number 2). The grains are not altered for depth of about 200 μm and 100 μm for alumina and zirconia blasted samples, respectively, where microstructure of the bulk is observed. Grain pattern of the bulk (bottom part) is characterized by a larger grain size (~ 25 μm) and the presence of wider thermally induced twins, obviously produced during the thermomechanical processing of the bar. Occasionally, elongated inclusions of MnS that are parallel to the longitudinal bar direction are observed. With regard to the effects of annealing on the microstructure of blasted specimens it is worth to remind that the recrystallization temperature is typically 1/3 to 1/2 of the melting temperature (can be as high as 0.7 Tm in some alloys). The recrystallization temperature, however, decreases as the cold work is increased, although below a “critical deformation” recrystallization does not occur. Taking into consideration these issues and the gradient in the plastic deformation during blasting, a non-homogeneous evolution of the microstructure during annealing is observed. After 2 min of annealing, Fig. 3, the nanograin size pattern is now more evident, which is consistent with an initial recovery step that allows a better crystallographic contrast of the image. After 1 h of annealing, signs of recrystallization are evident at the outermost part of the blasted affected zones, where the higher amount of cold work caused a larger decrease in the recrystallization temperature. As can be seen in Fig. 4, recrystallization causes the formation of new grains and a further growth achieving up to 1 μm in size, whereas the microstructure of the inner and less strained zone remains stable.
Fig. 1. Surface SEM images of specimens blasted with a) alumina and b) zirconia particles.
3.2. Magnetic force microscopy (MFM) thermal treatment, and especially after 1 h exposure, additional submicrometric particles, likely Fe- and Cr-rich oxides, which are not contributed to a roughness increase, are observed. Cross sectional examination, Fig. 2, reveals a grain size refinement beneath the blasted surface (see insets number 1), as expected [6,7]. Essentially, grain size ranges between 100 and 500 nm in a layer of about 30 μm for the BL-AlO samples, Fig. 2a, and 10–15 μm for the BL-ZrO ones, Fig. 2b. In the case of the specimens blasted with alumina, Fig. 2a,
XRD was used for phase identification but martensite peaks were not detected. MFM, which is more sensitive, revealed this phase and also allowed to study its distribution in the near surface region. A zone of about 40 μm in depth from the blasted surface was investigated. In the MFM images, dark and bright contrasts appear in certain regions of the samples. Dark contrast corresponds to attractive tip-sample interactions, while bright contrast can be associated with the regions where there is a repulsive tip-sample interaction.
Fig. 2. Cross sectional views of blasting affected zones corresponding to specimens blasted with a) alumina and b) zirconia particles. Insets (1) denote zones close to the blasted surfaces and inset (2) zones beneath the fine grained zone. SMn corresponds to inclusions of manganese sulphide.
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Fig. 3. Cross sectional views of blasting affected zones after thermal treatment at 700 °C for 2 min. a) Alumina blasted surfaces and b) zirconia blasted surfaces. Insets are close-up of zones beneath the blasted surfaces.
Fig. 4. Cross sectional views of blasting affected zones after thermal treatment at 700 °C for 1 H. a) Alumina blasted surfaces and b) zirconia blasted surfaces. Insets are close-up of zones beneath the blasted surfaces.
Fig. 5 shows typical topographic (Fig. 5a, b, c) and magnetic images (Fig. 5d, e, f) obtained on the cross sections, near to the surface, of the non-annealed (Fig. 5a, d) and annealed (Fig. 5b, c, e, f) zirconia blasted specimens. Comparison of Fig. 5a and d, corresponding to the same zone of the non-annealed and blasted specimen, indicates the lack of correlation between topography and magnetic signal, i.e. we can assume that the polishing process is not responsible for the magnetic behaviour. For the as-blasted condition, Fig. 5d, submicrometric α′martensite coexists with austenite characterized by the absence of magnetic signal. α′-Martensite is heterogeneously distributed along parallel and perpendicular directions to the surface. As we get closer to the surface, higher density of the magnetic phase is observed. Along the parallel direction, areas of high density of magnetic phase intercalate with areas of a lower density. In the transition area from the ultrafine to the non-affected microstructure typical α′-martensite needles can be found. Such lines present bright and dark contrast that corresponds to multidomain regions. In addition, dark spots that correspond to single domain crystals can also be found. Some “magnetic zones” present single domain configuration with low coercive fields, i.e. the magnetization in these regions will be always oriented parallel to the tip stray field, presenting homogeneous dark contrast. Thermal treatment of the BL_ZrO samples considerably reduces the amount of α′martensite and after 2 min, Fig. 5e, only some magnetic signal is observed at 15–20 μm from the surface. After 1 h of exposure, Fig. 5f, no magnetic signal is detected. In the BL_AlO sample, the MFM analysis reveals a more heterogeneous distribution of martensite along the blasted surface, with magnetic contrast at deeper regions, even at ~20 μm from the surface (images not shown). It is worth noting that the cumulative impact effect is necessary to produce α′-martensite formation, since a given threshold
strain level is needed for the transformation. The largest Al2O3 particles are more energetic than the ZrO2 particles and, therefore, lead to more cumulative deformation, which implies the formation of α′-martensite at deeper regions from the surface. On the other hand, their angular shape made that, depending of the attack angle to the surface, impact energy of some of them is devoted to tearing off material, leads to a micro-cutting effect, rather than to accumulate deformation. Hence, some of the α′-martensite accumulated from the previous impacts (which do not erode material) is also eliminated from the surface. Obviously, this increasing transformation process is retarded when local material removal by the large and edge-like Al2O3 particles occurs. Evolution of martensite with increasing exposure time is similar to that described for the BL-ZrO samples. Thermal treatment is very effective in causing a reversion of the martensite, even after only 2 min of exposure. This situation is illustrated in Fig. 6 that shows very small zones with magnetic contrast, Fig. 6b, around a large embedded alumina particle. 3.3. Microhardness Microhardness measurements of the blasted specimens, Fig. 7, reveal a gradient in hardness perpendicular to the blasted surface with a maximum of ~ 360 HV0.01 close to the surface, irrespectively the particle used for blasting and the thermal condition. In all cases, hardness achieves a near constant value of about 180 HV0.01 at a depth of ~ 200 μm and ~ 100 μm into the bulk, depending on whether the blasting was performed with particles of alumina, Fig. 7a, or zirconia, Fig. 7b, respectively. Thermal treatments at 700 °C do not significantly influence the hardness gradients, although depending on the relative depth to the
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Fig. 5. Topographic (a, b, c) and magnetic (d, e, f) images obtained by AFM (a, b, c) and MFM (d, e, f) in the cross sections of BL_ZrO samples in the as-blasted condition (a, d) and after annealing at 700 °C for 2 min (b, e) and 1 h (c, f). The line mark at the right indicates the edge of the sample (blasted surface).
a)
b)
Al2O3
around 100 μm, for the zirconia blasted specimens, are irrelevant, which is consistent with the thermal stability of the bulk microstructure at this temperature. 3.4. Rotating bending fatigue tests In a previous work [4] it has been shown that fatigue resistance of the 316 LVM steel (~400 MPa) is very close to the yield strength value (430 MPa). Interestingly, whereas blasting with zirconia rounded
Microhardness HV 0.01
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400
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BL_ZrO
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Fig. 6. Topographic (a) and magnetic (b) images obtained by AFM (a) and MFM (b) in the cross sections of BL_AlO samples after annealing at 700 °C for 1 h. The line mark at the right indicates the edge of the sample (blasted surface) and the circles delimit zones of magnetic signal.
BL_AlO
350
150
Microhardness HV 0.01
blasted surface slight changes in hardness are observed. Fig. 7 shows hardness values of blasted specimens with and without heat treatment at depths of about 10, 30, 100 and 200 μm from the blasted surface. The differences in hardness with regard to the bulk are likely related to the presence of α′-martensite, which has higher hardness than the austenite [23], the grain refinement, as the Hall–Petch expression predicts, and work hardening. As can be seen, thermal treatment for 2 min causes a moderated decrease in hardness. Despite the fact that both samples have approximately the same quantity of α′-martensite phase, the decrease in hardness for the alumina blasted specimens is more relevant at around 100 μm from the surface, which is consistent with the presence of less areas showing magnetic contrast close to the surface (Fig. 6) and a lesser dislocation density. Further increase in hardness with increasing exposure for a given depth can be associated with the grain size refinement associated with the reversion of the α′-martensite. According to Lee et al. [24] during isothermal holding, the equiaxed grains are nucleated and grown due to the diffusive reverse transformation, resulting in equiaxed grains, additional austenite volume fraction, and compressive transformation strain. It is worth to notice that grain growth observed at the outermost part of the blasted affected zones, Fig. 4a and b, is out the zone of measurements. Variations in hardness at around 200 μm in depth, for the alumina blasted specimens, and at
10
30 100 Depth (µm)
200
Fig. 7. Subsurface hardness as a function of depth from the surface for the specimens blasted with a) alumina and b) zirconia particles.
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particles causes a moderated increase in the fatigue limit (~420 MPa), blasting with angular alumina particles yields a significant decrease (~360 MPa). The higher roughness and the presence of a high amount of embedded alumina particles, acting as stress concentrators, make the alumina blasted surface prone to crack nucleation and growth despite the subsurface compressive residual stresses and grain size refinement. Thermal treatment of the blasted specimens at 700 °C for 2 min and 1 h influences the fatigue limit, as summarised in Table 1. In the case of the specimens blasted with zirconia particles, short term annealing hardly affects the fatigue strength. After 1 h of annealing, however, a moderated decrease in the fatigue strength (b5%) is observed, achieving values typical of the non-blasted condition. Interestingly, thermal treatment of the alumina blasted specimens causes an increase in the fatigue strength (~ 8%) after short annealing times (2 min), Fig. 8, achieving values for fatigue strength of about 400 MPa. The long annealing treatment, however, decreases the fatigue strength to minimum values (340 MPa). The beneficial effect of the short annealing treatment on the fatigue behaviour of the alumina blasted condition must be analysed by considering the microstructural induced changes influencing the crack initiation process [25]. Due to the very short annealing, the compressive residual stresses developed during blasting, which will serve to retard fatigue-crack propagation, are only partially released, as determined for similar samples by non-contacting thermoelectric measurements [20]. It is worth mentioning that partial relaxation of the compressive residual stresses would be counterbalanced by the reversion of α′-martensite and the induced grain size refinement of the microstructure [26], which would play a beneficial role. It could be argued that zones experiencing a plastic strain during fatigue, obviously at the outermost zone, could develop new strain induced martensite. Although cracks would develop around alumina particles acting as stress concentrators [24], the nanograin size of austenite is likely being controlling the damage process by acting as indirect barriers for crack propagation. The crack which meets the first grain boundary is shorter in the fine microstructure and would induce less strain at the crack tip, minimizing the formation of new martensite. A question still arises about why the beneficial effect of the short term annealing is not manifested for the zirconia blasted specimens. Blasted affected zone is larger for the alumina blasted specimens. Moreover, from the analysis of the gradient in hardness it follows that plastic deformation is more severe when blasting with alumina particles. Therefore, a larger amount of strain induced martensite is expected to form in the alumina samples. The detrimental role of the longest exposure can be associated with the full release of the initial compressive residual stresses, [20]. The lower fatigue strength of the thermally treated steel blasted with alumina particles can be related to their lower thermal expansion coefficient with regard to that of the matrix where they reside. During cooling from the annealing temperature, textural residual stresses become compressive in the ceramic particles but tensile in the matrix [27], favouring crack formation. From the above discussion it follows that short-term thermal treatments of grit blasted austenitic stainless steels may be a useful tool to reverse α′-martensite developed at the blasted affected zone as consequence of the severe plastic deformation, resulting in an increase in the fatigue strength. Improvement on the corrosion behaviour and ion Table 1 Rotating (R = −1) fatigue limit (MPa) of alumina and zirconia blasted specimens with and without post thermal treatments. Material condition
Non-treated
Polished BL-ZrO BL-AlO
400 420 370
Thermally treated (700 °C) 2 min
1h
– 420 390–400
– 390–400 340
600 550
Stress (MPa)
1080
500 450 400 350 300
SG-Alumina o SG-Alumina + 700 C/2min o SG-Alumina + 700 C/1h
250 200 4
5
10
10
6
10
7
10
8
10
Number of cycles to failure Fig. 8. Stress amplitude versus number of cycles to failure during rotating bending tests of alumina blasted steel before and after annealing for 2 min and 1 h.
release, which are dominated by the surface properties, cannot be discarded. Therefore, in vitro experiments considering short and medium periods of immersion in simulated human fluids are in progress. 4. Conclusions From the above discussion the following conclusions can be drawn: • Grit blasting of 316 LVM austenitic steel with either alumina or zirconia particles causes a severe surface plastic deformation that roughs the surface, decreases grain size, and favours the formation of martensite, which is in agreement with the previous works. • Annealing at 700 °C reverse the strain induced martensite (ferromagnetic) at the blasted affected zone to austenite (paramagnetic) without compromising the stability of the bulk microstructure, which eliminates the eventual problem associated with the use of strong magnetic fields during clinical diagnostics. • Blasting with rounded zirconia particles increases the fatigue resistance and short term annealing (2 min) hardly affects such increase. After 1 h of annealing, however, a moderated decrease in the fatigue strength (b5%) is observed, achieving values typical of the nonblasted condition. • Blasting with alumina particles, which are more abrasive and causes a huge amount of embedded particles, causes a significant decrease in the fatigue strength. Interestingly, short annealing times (2 min) cause an increase in the fatigue strength (~ 8%) as consequence of the reversion of the martensite, achieving values corresponding to the non-blasted condition. Longer periods of annealing cause a moderate grain size increase at the outermost surface and contribute to full relaxation of the compressive residual stresses, yielding a significant reduction in the fatigue strength. Conflict of interest None. Acknowledgements The authors are thankful for financial support from the MICINN (MAT2009-14695-C04) and the MINECO (MAT 2012-37736-C05-01). S. Barriuso thanks to JAE predoc grant of CSIC. Technicians of Laboratories of X-Ray Diffraction, Mechanical Testing, and Microscopy (CENIM) and the SPM-Service of the ICMM-CSIC are specially acknowledged. References [1] J.A. Disegi, L. Eschbach, Injury 31 (Suppl. 4) (2000) 2–6.
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