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Mechanical and thermal stability of mechanically induced near-surface nanostructures
Materials Science and Engineering A 403 (2005) 318–327
Mechanical and thermal stability of mechanically induced near-surface nanostructures I. Nikitin a , I. Altenberger a,∗ , H.J. Maier b , B. Scholtes a b
a Institute of Materials Engineering, University of Kassel, Germany Lehrstuhl f¨ur Werkstoffkunde (Materials Science), University of Paderborn, Germany
Received in revised form 9 May 2005; accepted 13 May 2005
Abstract Mechanical surface treatments, such as deep rolling, shot peening, hammering, etc., can significantly improve the fatigue behaviour of metallic materials owing to near-surface nanocrystallisation, strain hardening and compressive residual stresses. In this paper, we investigate the stability of near-surface microstructures of deep rolled austenitic stainless steel AISI 304 and turbine blade alloy Ti–6Al–4V during high temperature fatigue (up to 600 ◦ C) by transmission electron microscopy and X-ray diffraction. The investigated nanocrystalline regions are stable during short time annealing and unstable during long time annealing at 600 ◦ C. Isothermal fatigue in the low cycle fatigue regime at high stress amplitudes does not alter the nanocrystalline region up to 600 ◦ C. © 2005 Elsevier B.V. All rights reserved. Keywords: Surface nanocrystallisation; Nanocrystalline materials; Transmission electron microscopy; Microstructure; Fatigue; Deep rolling
1. Introduction Nanocrystalline metallic materials have been reported to offer a variety of beneficial properties, such as, e.g. higher biocompatibility [1], excellent magnetic properties [2], higher hardness [3,4] and wear resistance [5–7], as well as higher resistance against fatigue crack initiation [8] as compared to non-nanocrystalline conventional materials. The production of nanocrystalline bulk materials is up to now expensive and very time-, equipment- and cost-intensive. However, most failures of engineering materials, such as, e.g. fatigue crack nucleation, corrosion or wear take place at the surface. Therefore, alternatives to bulk nanocrystalline materials can be materials with a nanocrystalline surface layer, which is easier to produce and more cost-efficient. Nanocrystalline layers can be produced by basic mechanical surface treatments, such as shot peening [8,9], deep rolling [10], hammering, etc. [7,11]. The stability of these nearsurface regions against mechanical and thermal loading, especially under fatigue conditions, is a key issue, and the ∗
Corresponding author. Tel.: +49 561 804 3678; fax: +49 561 804 3662. E-mail address: [email protected] (I. Altenberger).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.05.030
present study addresses the stability of the nanocrystalline microstructure produced by deep rolling. The basic principles of mechanical surface treatments are well known: a localised elastic–plastic deformation in nearsurface regions leads to the formation of compressive residual stresses and severe microstructural alterations, enabling the strengthened near-surface regions to exhibit higher resistance against fatigue crack initiation and propagation [12–15]. Deep rolling improves the fatigue behaviour over a wide temperature range up to 600 ◦ C owing to near-surface nanocrystallisation, strain hardening and compressive residual stresses [16–18]. Moreover, metastable austenitic steels, such as AISI 304, exhibit deformation-induced phase transformations. The observed microstructural alterations in near-surfaces regions for the investigated alloys are complex. Basically, two layers or regions can be distinguished: a nanocrystalline layer directly at the surface and a layer exhibiting very high dislocation densities immediately below that surface layer. For AISI 304, compressive residual stresses start to relax already during fatigue at 100 ◦ C if the stress amplitude is in the low cycle fatigue regime [19]. The improved fatigue life at elevated temperatures is primarily caused by a combination
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of strain hardening, strain induced martensitic transformation and a nanocrystalline layer [8,20]. 2. Materials and experimental methods 2.1. Materials The investigated materials were hot rolled cylindrical bars (14 mm diameter) of AISI 304 austenitic steel and turbine blade alloy Ti–6Al–4V. The microstructure of AISI 304 was fully austenitic with an average grain size of 70 m. Uniaxial tensile testing at 25 ◦ C gave a yield strength of 345 MPa, an ultimate tensile strength of 624 MPa, and 52% elongation to fracture. The Ti–6Al–4V alloy was solution-treated for 1 h at 970 ◦ C, air cooled, machined and then stress relieved for 2 h at 700 ◦ C. The resulting microstructure consisted of a distribution of interconnected equiaxed primary-␣ grains with a grain size of 10 m and lamellar ␣ +  colonies (transformed). Uniaxial tensile testing at 25 ◦ C gave yield strength of 979 MPa, an ultimate tensile strength of 1172 MPa, and 17% elongation to fracture. 2.2. Experimental methods Unnotched cylindrical specimens with a gauge length of 10 mm and a diameter of 7 mm were used for the surface treatment and subsequent fatigue studies. For deep rolling, a standard spherical rolling element (6.6 mm diameter) was used with a constant feed of 0.1 mm per revolution and a rolling pressure of 150 bar (rolling force 0.5 kN). Isothermal fatigue experiments were performed under load control in tension-compression using a standard servohydraulical testing machine with a load ratio of R = −1, a frequency of 5 Hz and temperatures of 25–600 ◦ C. The samples were heated by radiant heating. In order to minimise thermal gradients,
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heating of the smooth specimens was started 10 min prior to the start of the actual fatigue tests. The near-surface work hardening was characterised by full width at half maximum values (FWHM) of X-ray interference lines. In case of AISI 304 samples, Cr K␣-radiation was applied and the {2 2 0}Bragg peak of the austenite phase was analysed, whereas Cu K␣-radiation and the {2 1 3}-Bragg peak of the ␣-phase were used for the Ti–6Al–4V material. Transmission electron microscopy (TEM) was performed using a 200 kV Philips microscope. TEM foils from a given depth below the surface were prepared by a combination of single and twinjet electropolishing. Whenever possible, bright-field images were recorded under two-beam conditions. The stress axis is always oriented parallel to the TEM foils.
3. Results 3.1. Microstructure of the deep rolled state The metastable austenitic steel AISI 304 exhibits deformation-induced phase transformations and a high density of mechanical twins [11,18]. The calculation of the Md30 temperature, i.e. the temperature where 50% martensite is formed after 30% plastic deformation, is between 23 and –12 ◦ C for the investigated chemical composition [21]. The surface exhibits a roughly 2–3 m thick nanocrystalline layer (Fig. 1, left). It contains martensitic and austenitic crystallites with a grain size of about 30 nm. The thickness of the nanocrystalline layer is strongly dependent on the deep rolling parameters (rolling pressure and coverage). For example, for AISI 304 an increase of the rolling coverage by a factor of 10 increases the thickness of the nanocrystalline layer from 2 to about 6 m. The grain size in the nanocrystalline layer was about 30 nm directly at the surface and increased continuously to 40 nm in a depth of 1.4 m. The subsurface
Fig. 1. TEM micrographs showing the surface (left) and the subsurface region at a depth of roughly 5 m (right) of AISI 304.
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Fig. 2. TEM micrographs of the direct surface region of the deep rolled stainless steel AISI 304.
exhibits a layer with very high dislocation densities in the austenitic matrix between martensitic needles and mechanical twins (Fig. 1, right). The martensite content shows a subsurface maximum of about 5% and after that decreases with increasing distance from the surface. Moreover, similar to the martensite content, the concentration of mechanical twins also decreases with increasing distance from the surface. Fig. 2 shows the martensitic and austenitic nanocrystallites in higher resolution. It can be seen that the nanocrystallites have a complex lamellar substructure. In contrast to nanocrystallites produced by sandblasting [7] the nano-grains after deep rolling emerged without annealing after the mechanical treatment. In the as-received condition, the investigated turbine blade alloy Ti–6Al–4V has an average ␣-grain size of about 10 m. After deep rolling the surface exhibits a
nanocrystalline layer with crystallite sizes of about 50 nm directly at the surface (Fig. 3, left). The subsurface exhibits high dislocation densities (Fig. 3, right). Similar to AISI 304 the nanocrystallites have a complex lamellar substructure. 3.2. Microstructural stability during cyclic loading Microstructural stability during cyclic loading of nearsurface regions at room temperature is very crucial for the effectiveness of surface treatments, especially during fatigue loading. Figs. 4 and 5 show examples of microstructures in the direct surface region and for a region immediately beneath the nanocrystalline layer (at 5 m depth), respectively, after room temperature fatigue. The samples were fatigued at room temperature with a stress amplitude of σ a = 350 MPa. At this
Fig. 3. TEM micrographs of deep rolled Ti–6Al–4V directly at the surface (left) and in the subsurface region in a depth of roughly 5 m (right).
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Fig. 4. TEM micrograph of the direct surface region of the deep rolled fatigued surface of AISI 304 (σ a = 350 MPa, T = 25 ◦ C, after 1 (left) and 8000 cycles (right)).
stress amplitude, fracture occurred after 8600 cycles. TEM foils were prepared after 1, 10, 100, 1000 and 8000 cycles and TEM-images under two-beam conditions were recorded. Fig. 4 demonstrates that the microstructure of the nanocrystalline layer remains stable under these loading conditions, i.e. the nanocrystalline layer before and after fatigue is almost identical. Initially, the subsurface in a depth of 5 m shows very high dislocation densities between martensitic needles and mechanical twins (Fig. 5, left). During cyclic loading the dislocation density in the austenite phase decreases continuously. The decrease of the dislocation density is also reflected in the decline of FWHM values of X-ray diffraction peaks with increasing numbers of cycles. For the assessment of the measured FWHM values it should be pointed
out that the FWHM value depends on several factors, such as dislocation densities and line broadening by crystal size effects. In the case investigated, the FWHM value is an integral average value representing diffracted information from a surface region up to 5 m in a depth. The separation between crystal size and broadening from strain fields was not carried out in the case investigated. After 8000 cycles the FWHM value in the subsurface region has decreased by more than 40% compared to the initial state after the surface treatment. It is evident that cyclic loading did not affect the nanocrystalline layer significantly. However, the subsurface non-nanocrystalline layer with high initial dislocation densities is strongly affected by mechanical fatigue. The dislocation density decreases both with increasing distance from
Fig. 5. TEM micrographs of the subsurface region (5 m) of the deep rolled AISI 304 (σ a = 350 MPa, T = 25 ◦ C, for 1 (left) and 8000 cycles (right)).
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the surface and with increasing numbers of cycles. Obviously, the microstructure of the subsurface layer is not stable during mechanical fatigue. Fatigue-induced dislocation cell structures as reported for shot peened specimens cycled at high stress amplitudes [8] were not detected in the present study in the regions investigated. 3.3. Microstructural stability at elevated temperatures Different annealing treatments were applied to shed light on the thermal stability of the nanocrystalline surface layer. TEM investigations indicated that in deep rolled AISI 304 the microstructure in this layer is stable during short isothermal annealing up to 600 ◦ C for up to 0.3 h annealing time. However, annealing at 600 ◦ C for 1 h resulted in an increase of the grain size in the nanocrystalline layer to about 70 nm (Fig. 6). Further annealing at 600 ◦ C for 10 and 100 h lead to further grain growth. The grain size after annealing at 600 ◦ C for 10 h is about 150–200 nm. After 100 h annealing at 600 ◦ C the grain size exceeded 200 nm. Among others, the grain growth of the nanocrystals is reflected by the decrease of the FWHM value of X-ray diffraction peaks (Fig. 7). The microstructure of the subsurface layer exhibiting a high dislocation density is unstable and changes during thermal exposure as well. The decrease of dislocation densities is reflected indirectly in the decrease of the FWHM value (Fig. 7). Figs. 7 and 8 give an impression of the change of the microstructures in the subsurface layer. Fig. 7 reveals that up to 450 ◦ C the annealed state (annealing time 1 h) shows almost the same work hardening state as the as deep rolled condition. By contrast, annealing at 600 ◦ C leads to a marked decrease of FWHM values even at short annealing times. At the surface, the FWHM values decreased by as much as 32% as compared to the as deep rolled condition. Whereas an increase of the annealing temperature had a marked effect, an increase of annealing time at 600 ◦ C from 10 to 100 h did not change the FWHM values
Fig. 7. Temperature- and time-dependence of the FWHM value of X-ray diffraction peaks of deep rolled AISI 304 during annealing as a function of distance to surface.
drastically. For both annealing times it decreased by about 50% as compared to the as deep rolled state. 3.4. Stability against cyclic loading at elevated temperature Stability against isothermal cyclic loading of the nearsurface nanocrystallites in AISI 304 was analysed in isothermal stress controlled fatigue tests at elevated temperatures with stress amplitude of 280 MPa. These tests were interrupted after 2000 cycles. For the test frequency of 5 Hz, 2000 cycles correspond to 7 min testing time at elevated temperature plus 10 min for the pre-heating to reach thermal equilibrium prior to the test. TEM observations demonstrated that the nanocrystalline layer was also stable up to 600 ◦ C under these conditions (Fig. 9). A quantitative determination of the grain size revealed that the grain size within the whole nanocrystalline layer had not changed significantly after high temperature fatigue (T = 450 ◦ C, σ a = 280 MPa, N = Nf /2) as compared to the as deep rolled condition. The subsurface
Fig. 6. TEM micrographs of the deep rolled and annealed surface region of AISI 304 for different annealing times at 600 ◦ C (1 h (left) and 100 h (right)).
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Fig. 8. TEM micrograph of the deep rolled and annealed subsurface of AISI 304 for different annealing times at 600 ◦ C (1 h (left) and 100 h (right)).
Fig. 9. TEM micrographs of the deep rolled and fatigued surface of AISI 304 (σ a = 280 MPa, 2000 cycles, 100 ◦ C (left) and 600 ◦ C (right)).
shows a decrease of dislocation density with increasing test temperature, the FWHM value in a depth of 0.01 mm dropped by 41% after isothermal fatigue at 600 ◦ C as compared to the as deep rolled condition (Figs. 10 and 11). The thermomechanical stability of the nanocrystalline surface in Ti–6Al–4V was determined in fatigue tests conducted at different temperatures and a constant stress amplitude of 430 MPa. Similar to AISI 304, it was observed that the nanocrystalline layer is stable up to 600 ◦ C under these conditions (Fig. 12). Moreover, the subsurface (5 m depth) shows a decrease of dislocation density with increasing temperature (Fig. 13). The microstructural changes in deep rolled AISI 304 as a function of number of cycles for a test temperature of 500 ◦ C and a stress amplitude of σ a = 280 MPa are shown in Figs. 14 and 15. Fracture in this case occurred after 4500
Fig. 10. Influence of test temperature on the half-width depth profile (σ a = 280 MPa, 2000 cycles).
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Fig. 11. TEM micrographs of the deep rolled and fatigued subsurface region of AISI 304 (5 m depth, σ a = 280 MPa, 2000 cycles, 100 ◦ C (left) and 600 ◦ C (right)).
cycles. The samples were fatigued and investigated after fixed loading cycles (1, 100, 1000 and 4000 cycles). Fig. 14 shows in an exemplary manner characteristic microstructures of the surface region and Fig. 15 of the subsurface region (5 m depth) for 1 and 4000 cycles corresponding to 0.02 and 88% of the total lifetime of the specimens. Again, the nanocrystalline layer is stable and the subsurface region is characterised by microstructural alterations and a decrease of dislocation density. At 4000 cycles (88% of fatigue lifetime) the subsurface showed tangled and dense dislocation arrangements. Diffuse cell structures parallel to the stress axis started to form (Fig. 15). In a previous study, the formation of a dislocation cell structure was observed in a deep rolled AISI 304 during cyclic deformation at high stress amplitudes [8]. It
can be summarised that isothermal fatigue up to 600 ◦ C (total time at elevated temperature of 17 min) does not affect the microstructure of the nanocrystalline layer, i.e. the shape and size of the grains both remained unchanged. The subsurface layer with its initially high dislocation density, however, is strongly affected and the dislocation density decreases during fatigue. As expected, the key parameter was the actual test temperature in these tests.
4. Discussion The fatigue life of metallic materials and components depends in a complex way on a multitude of factors among
Fig. 12. TEM micrographs of the deep rolled and fatigued surface region of Ti–6Al–4V (σ a = 430 MPa, 2000 cycles, 100 ◦ C (left) and 500 ◦ C (right)).
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Fig. 13. TEM micrograph of the deep rolled and fatigued subsurface of Ti–6Al–4V (σ a = 430 MPa, 2000 cycles, 100 ◦ C (left) and 500 ◦ C (right)).
which near-surface residual stresses and microstructures as well as roughness are just the most important ones for surface treated components. In assessing residual stress effects some considerable progress has been made over the past decades [19,22,25] and residual stress effects on fatigue performance have been thoroughly investigated even for complex parts and loading conditions [23,24,26]. Near-surface microstructures have been investigated much more rarely [9,15,19]. Moreover, surface treated components may experience alterations of near-surface properties during service load, especially under combined mechanical and thermal loading, such as in power plant or aircraft applications. The present study confirms that mechanical surface treatments, such as deep rolling, induce nanocrystalline
surfaces and highly strain hardened layers. The phenomenon of mechanically induced nanocrystallisation has also been observed in other contexts and related problems such as chip forming processes, hard turning or wear [29–31]. Nanocrystalline surface layers significantly affect the mechanical behaviour by restricting or impeding dislocation slip and the formation of slip bands at the surface [8], which act as preferential crack initiation sites. Therefore, taylored nanocrystalline surfaces are highly desirable in order to delay or prevent surface fatigue damage. One critical issue of mechanically induced nanocrystalline surface layers is their mechanical and thermal stability. Firstly, thermodynamically, nanocrystalline regions are in a non-equilibrium state with a high number of lattice defects that tend to anneal out, and finally recrystallise if
Fig. 14. TEM micrograph of the deep rolled and fatigued surface of AISI 304 (σ a = 280 MPa, 500 ◦ C, after own cycle (left) and 4000 cycles (right)).
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Fig. 15. TEM micrograph of the deep rolled and fatigued subsurface of AISI 304 (σ a = 280 MPa, 500 ◦ C, for 1 (left) and 4000 cycles (right)).
the temperatures are elevated enough to provide substantial diffusion processes. Secondly, according to recent in situ TEM studies [28], dislocation slip occurs in nanocrystalline materials, if the grain size is in the order of 20–50 nm or higher. Thus, it appeared therefore probable, that mechanical (cyclic or monotonic) loading affects the long-term or even short-term stability of nanocrystalline surface regions. By contrast, the present results indicate that nanocrystalline surface layers in AISI 304 and Ti–6Al–4V induced by deep rolling exhibit an extraordinarily high mechanical stability under the given conditions and remain very stable at high cyclic stress or plastic strain in the low-cycle fatigue (LCF) regime. This implies, that also in the high-cycle fatigue (HCF) regime where the plastic strain is much lower, mechanical stability of nanocrystalline surface layers is to be expected at low-test temperatures. The isothermal annealing experiments revealed that considerable grain coarsening takes place with increasing temperature and annealing time. In situ heating of TEM foils have yielded information on the thermal stability of nanocrystalline regions under shortterm annealing conditions [27]. An annealing temperature of 600 ◦ C can render the nanocrystalline surface layers “microcrystalline” after many hours of annealing as indicated by the present studies. Interestingly, also isothermal fatigue at elevated temperatures in the LCF regime (short LCF test) did not alter the structure of the nanocrystalline surface regions significantly in spite of simultaneous elevated temperature exposure and high cyclic plastic strain amplitude (which was in the order of 0.1–0.5%). Obviously, in the short LCF tests the exposure time is not sufficient to permit substantial annealing processes, therefore preventing coarsening and recrystallisation. We note that the results obtained from the presented investigations are not to be generalised for other test frequencies. The stability of nanocrystalline layers against high
temperature fatigue is clearly affected much more by the exposure time than by the number of cycles. Lower frequencies will inevitably lead to longer exposure time, and thus, to more pronounced diffusion processes and higher instability of the nanocrystallites. As a second consequence, instability of nanocrystalline layers may occur in the HCF regime, even though stability occurs in the LCF regime, because of different exposure times inspite of identical test temperatures. The more critical loading case may thus be the HCF case and not the LCF case. For high temperature fatigue it appears that the stability of nanocrystalline surface layers is mostly influenced by the temperature, followed by the exposure time, and finally by the stress/strain amplitude and number of cycles. Finally, the present study shows that the stability of nearsurface microstructures also strongly depends on the nature of near-surface microstructures: Underneath the nanocrystalline layer, the dislocation arrangements are generally much more instable to high temperature fatigue and transform into annealed low-energy arrangements with increasing temperature and increasing numbers of cycles as also indicated indirectly by FWHM value measurements.
5. Summary The mechanical, thermal and thermomechanical stability of nanostructured surface regions induced by deep rolling has been investigated in AISI 304 type metastable austenitic stainless steel and fan blade alloy Ti–6Al–4V. The results can be summarised as follows. • Deep rolling induced a nanocrystalline layer in both materials, deformation-induced martensite and mechanical twins in AISI 304 as well as high macro and micro residual stresses due to very high dislocation densities [19].
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• The nanocrystalline layer is stable during short time annealing up to 600 ◦ C. For AISI 304 long time annealing at 600 ◦ C (10–100 h) leads to recrystallisation. After 10 h at 600 ◦ C the nanocrystalline layer transforms into a fine equiaxed grain structure with an average grain size of about 150–200 nm. After 100 h at 600 ◦ C almost complete recrystallisation of the nanocrystalline layer takes place. • Isothermal fatigue in the LCF regime at high stress amplitudes does not significantly alter the structure of the nanocrystalline surface layer in the temperature range 25–600 ◦ C for the chosen test conditions. • Subsurface non-nanocrystalline regions with high dislocation densities tend to change into low energy dislocation arrangements during thermal, mechanical or thermomechanical loading. In all cases, a decline of the FWHM values of diffraction peaks during loading was observed.
Acknowledgement The authors would like to thank the German Science Foundation (DFG) for financial support of the Emmy-Noether group in Kassel (under contract-number AL 558/1-2).
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Mechanical and thermal stability of mechanically induced near-surface nanostructures I. Nikitin a , I. Altenberger a,∗ , H.J. Maier b , B. Scholtes a b
a Institute of Materials Engineering, University of Kassel, Germany Lehrstuhl f¨ur Werkstoffkunde (Materials Science), University of Paderborn, Germany
Received in revised form 9 May 2005; accepted 13 May 2005
Abstract Mechanical surface treatments, such as deep rolling, shot peening, hammering, etc., can significantly improve the fatigue behaviour of metallic materials owing to near-surface nanocrystallisation, strain hardening and compressive residual stresses. In this paper, we investigate the stability of near-surface microstructures of deep rolled austenitic stainless steel AISI 304 and turbine blade alloy Ti–6Al–4V during high temperature fatigue (up to 600 ◦ C) by transmission electron microscopy and X-ray diffraction. The investigated nanocrystalline regions are stable during short time annealing and unstable during long time annealing at 600 ◦ C. Isothermal fatigue in the low cycle fatigue regime at high stress amplitudes does not alter the nanocrystalline region up to 600 ◦ C. © 2005 Elsevier B.V. All rights reserved. Keywords: Surface nanocrystallisation; Nanocrystalline materials; Transmission electron microscopy; Microstructure; Fatigue; Deep rolling
1. Introduction Nanocrystalline metallic materials have been reported to offer a variety of beneficial properties, such as, e.g. higher biocompatibility [1], excellent magnetic properties [2], higher hardness [3,4] and wear resistance [5–7], as well as higher resistance against fatigue crack initiation [8] as compared to non-nanocrystalline conventional materials. The production of nanocrystalline bulk materials is up to now expensive and very time-, equipment- and cost-intensive. However, most failures of engineering materials, such as, e.g. fatigue crack nucleation, corrosion or wear take place at the surface. Therefore, alternatives to bulk nanocrystalline materials can be materials with a nanocrystalline surface layer, which is easier to produce and more cost-efficient. Nanocrystalline layers can be produced by basic mechanical surface treatments, such as shot peening [8,9], deep rolling [10], hammering, etc. [7,11]. The stability of these nearsurface regions against mechanical and thermal loading, especially under fatigue conditions, is a key issue, and the ∗
Corresponding author. Tel.: +49 561 804 3678; fax: +49 561 804 3662. E-mail address: [email protected] (I. Altenberger).
0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.05.030
present study addresses the stability of the nanocrystalline microstructure produced by deep rolling. The basic principles of mechanical surface treatments are well known: a localised elastic–plastic deformation in nearsurface regions leads to the formation of compressive residual stresses and severe microstructural alterations, enabling the strengthened near-surface regions to exhibit higher resistance against fatigue crack initiation and propagation [12–15]. Deep rolling improves the fatigue behaviour over a wide temperature range up to 600 ◦ C owing to near-surface nanocrystallisation, strain hardening and compressive residual stresses [16–18]. Moreover, metastable austenitic steels, such as AISI 304, exhibit deformation-induced phase transformations. The observed microstructural alterations in near-surfaces regions for the investigated alloys are complex. Basically, two layers or regions can be distinguished: a nanocrystalline layer directly at the surface and a layer exhibiting very high dislocation densities immediately below that surface layer. For AISI 304, compressive residual stresses start to relax already during fatigue at 100 ◦ C if the stress amplitude is in the low cycle fatigue regime [19]. The improved fatigue life at elevated temperatures is primarily caused by a combination
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of strain hardening, strain induced martensitic transformation and a nanocrystalline layer [8,20]. 2. Materials and experimental methods 2.1. Materials The investigated materials were hot rolled cylindrical bars (14 mm diameter) of AISI 304 austenitic steel and turbine blade alloy Ti–6Al–4V. The microstructure of AISI 304 was fully austenitic with an average grain size of 70 m. Uniaxial tensile testing at 25 ◦ C gave a yield strength of 345 MPa, an ultimate tensile strength of 624 MPa, and 52% elongation to fracture. The Ti–6Al–4V alloy was solution-treated for 1 h at 970 ◦ C, air cooled, machined and then stress relieved for 2 h at 700 ◦ C. The resulting microstructure consisted of a distribution of interconnected equiaxed primary-␣ grains with a grain size of 10 m and lamellar ␣ +  colonies (transformed). Uniaxial tensile testing at 25 ◦ C gave yield strength of 979 MPa, an ultimate tensile strength of 1172 MPa, and 17% elongation to fracture. 2.2. Experimental methods Unnotched cylindrical specimens with a gauge length of 10 mm and a diameter of 7 mm were used for the surface treatment and subsequent fatigue studies. For deep rolling, a standard spherical rolling element (6.6 mm diameter) was used with a constant feed of 0.1 mm per revolution and a rolling pressure of 150 bar (rolling force 0.5 kN). Isothermal fatigue experiments were performed under load control in tension-compression using a standard servohydraulical testing machine with a load ratio of R = −1, a frequency of 5 Hz and temperatures of 25–600 ◦ C. The samples were heated by radiant heating. In order to minimise thermal gradients,
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heating of the smooth specimens was started 10 min prior to the start of the actual fatigue tests. The near-surface work hardening was characterised by full width at half maximum values (FWHM) of X-ray interference lines. In case of AISI 304 samples, Cr K␣-radiation was applied and the {2 2 0}Bragg peak of the austenite phase was analysed, whereas Cu K␣-radiation and the {2 1 3}-Bragg peak of the ␣-phase were used for the Ti–6Al–4V material. Transmission electron microscopy (TEM) was performed using a 200 kV Philips microscope. TEM foils from a given depth below the surface were prepared by a combination of single and twinjet electropolishing. Whenever possible, bright-field images were recorded under two-beam conditions. The stress axis is always oriented parallel to the TEM foils.
3. Results 3.1. Microstructure of the deep rolled state The metastable austenitic steel AISI 304 exhibits deformation-induced phase transformations and a high density of mechanical twins [11,18]. The calculation of the Md30 temperature, i.e. the temperature where 50% martensite is formed after 30% plastic deformation, is between 23 and –12 ◦ C for the investigated chemical composition [21]. The surface exhibits a roughly 2–3 m thick nanocrystalline layer (Fig. 1, left). It contains martensitic and austenitic crystallites with a grain size of about 30 nm. The thickness of the nanocrystalline layer is strongly dependent on the deep rolling parameters (rolling pressure and coverage). For example, for AISI 304 an increase of the rolling coverage by a factor of 10 increases the thickness of the nanocrystalline layer from 2 to about 6 m. The grain size in the nanocrystalline layer was about 30 nm directly at the surface and increased continuously to 40 nm in a depth of 1.4 m. The subsurface
Fig. 1. TEM micrographs showing the surface (left) and the subsurface region at a depth of roughly 5 m (right) of AISI 304.
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Fig. 2. TEM micrographs of the direct surface region of the deep rolled stainless steel AISI 304.
exhibits a layer with very high dislocation densities in the austenitic matrix between martensitic needles and mechanical twins (Fig. 1, right). The martensite content shows a subsurface maximum of about 5% and after that decreases with increasing distance from the surface. Moreover, similar to the martensite content, the concentration of mechanical twins also decreases with increasing distance from the surface. Fig. 2 shows the martensitic and austenitic nanocrystallites in higher resolution. It can be seen that the nanocrystallites have a complex lamellar substructure. In contrast to nanocrystallites produced by sandblasting [7] the nano-grains after deep rolling emerged without annealing after the mechanical treatment. In the as-received condition, the investigated turbine blade alloy Ti–6Al–4V has an average ␣-grain size of about 10 m. After deep rolling the surface exhibits a
nanocrystalline layer with crystallite sizes of about 50 nm directly at the surface (Fig. 3, left). The subsurface exhibits high dislocation densities (Fig. 3, right). Similar to AISI 304 the nanocrystallites have a complex lamellar substructure. 3.2. Microstructural stability during cyclic loading Microstructural stability during cyclic loading of nearsurface regions at room temperature is very crucial for the effectiveness of surface treatments, especially during fatigue loading. Figs. 4 and 5 show examples of microstructures in the direct surface region and for a region immediately beneath the nanocrystalline layer (at 5 m depth), respectively, after room temperature fatigue. The samples were fatigued at room temperature with a stress amplitude of σ a = 350 MPa. At this
Fig. 3. TEM micrographs of deep rolled Ti–6Al–4V directly at the surface (left) and in the subsurface region in a depth of roughly 5 m (right).
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Fig. 4. TEM micrograph of the direct surface region of the deep rolled fatigued surface of AISI 304 (σ a = 350 MPa, T = 25 ◦ C, after 1 (left) and 8000 cycles (right)).
stress amplitude, fracture occurred after 8600 cycles. TEM foils were prepared after 1, 10, 100, 1000 and 8000 cycles and TEM-images under two-beam conditions were recorded. Fig. 4 demonstrates that the microstructure of the nanocrystalline layer remains stable under these loading conditions, i.e. the nanocrystalline layer before and after fatigue is almost identical. Initially, the subsurface in a depth of 5 m shows very high dislocation densities between martensitic needles and mechanical twins (Fig. 5, left). During cyclic loading the dislocation density in the austenite phase decreases continuously. The decrease of the dislocation density is also reflected in the decline of FWHM values of X-ray diffraction peaks with increasing numbers of cycles. For the assessment of the measured FWHM values it should be pointed
out that the FWHM value depends on several factors, such as dislocation densities and line broadening by crystal size effects. In the case investigated, the FWHM value is an integral average value representing diffracted information from a surface region up to 5 m in a depth. The separation between crystal size and broadening from strain fields was not carried out in the case investigated. After 8000 cycles the FWHM value in the subsurface region has decreased by more than 40% compared to the initial state after the surface treatment. It is evident that cyclic loading did not affect the nanocrystalline layer significantly. However, the subsurface non-nanocrystalline layer with high initial dislocation densities is strongly affected by mechanical fatigue. The dislocation density decreases both with increasing distance from
Fig. 5. TEM micrographs of the subsurface region (5 m) of the deep rolled AISI 304 (σ a = 350 MPa, T = 25 ◦ C, for 1 (left) and 8000 cycles (right)).
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the surface and with increasing numbers of cycles. Obviously, the microstructure of the subsurface layer is not stable during mechanical fatigue. Fatigue-induced dislocation cell structures as reported for shot peened specimens cycled at high stress amplitudes [8] were not detected in the present study in the regions investigated. 3.3. Microstructural stability at elevated temperatures Different annealing treatments were applied to shed light on the thermal stability of the nanocrystalline surface layer. TEM investigations indicated that in deep rolled AISI 304 the microstructure in this layer is stable during short isothermal annealing up to 600 ◦ C for up to 0.3 h annealing time. However, annealing at 600 ◦ C for 1 h resulted in an increase of the grain size in the nanocrystalline layer to about 70 nm (Fig. 6). Further annealing at 600 ◦ C for 10 and 100 h lead to further grain growth. The grain size after annealing at 600 ◦ C for 10 h is about 150–200 nm. After 100 h annealing at 600 ◦ C the grain size exceeded 200 nm. Among others, the grain growth of the nanocrystals is reflected by the decrease of the FWHM value of X-ray diffraction peaks (Fig. 7). The microstructure of the subsurface layer exhibiting a high dislocation density is unstable and changes during thermal exposure as well. The decrease of dislocation densities is reflected indirectly in the decrease of the FWHM value (Fig. 7). Figs. 7 and 8 give an impression of the change of the microstructures in the subsurface layer. Fig. 7 reveals that up to 450 ◦ C the annealed state (annealing time 1 h) shows almost the same work hardening state as the as deep rolled condition. By contrast, annealing at 600 ◦ C leads to a marked decrease of FWHM values even at short annealing times. At the surface, the FWHM values decreased by as much as 32% as compared to the as deep rolled condition. Whereas an increase of the annealing temperature had a marked effect, an increase of annealing time at 600 ◦ C from 10 to 100 h did not change the FWHM values
Fig. 7. Temperature- and time-dependence of the FWHM value of X-ray diffraction peaks of deep rolled AISI 304 during annealing as a function of distance to surface.
drastically. For both annealing times it decreased by about 50% as compared to the as deep rolled state. 3.4. Stability against cyclic loading at elevated temperature Stability against isothermal cyclic loading of the nearsurface nanocrystallites in AISI 304 was analysed in isothermal stress controlled fatigue tests at elevated temperatures with stress amplitude of 280 MPa. These tests were interrupted after 2000 cycles. For the test frequency of 5 Hz, 2000 cycles correspond to 7 min testing time at elevated temperature plus 10 min for the pre-heating to reach thermal equilibrium prior to the test. TEM observations demonstrated that the nanocrystalline layer was also stable up to 600 ◦ C under these conditions (Fig. 9). A quantitative determination of the grain size revealed that the grain size within the whole nanocrystalline layer had not changed significantly after high temperature fatigue (T = 450 ◦ C, σ a = 280 MPa, N = Nf /2) as compared to the as deep rolled condition. The subsurface
Fig. 6. TEM micrographs of the deep rolled and annealed surface region of AISI 304 for different annealing times at 600 ◦ C (1 h (left) and 100 h (right)).
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Fig. 8. TEM micrograph of the deep rolled and annealed subsurface of AISI 304 for different annealing times at 600 ◦ C (1 h (left) and 100 h (right)).
Fig. 9. TEM micrographs of the deep rolled and fatigued surface of AISI 304 (σ a = 280 MPa, 2000 cycles, 100 ◦ C (left) and 600 ◦ C (right)).
shows a decrease of dislocation density with increasing test temperature, the FWHM value in a depth of 0.01 mm dropped by 41% after isothermal fatigue at 600 ◦ C as compared to the as deep rolled condition (Figs. 10 and 11). The thermomechanical stability of the nanocrystalline surface in Ti–6Al–4V was determined in fatigue tests conducted at different temperatures and a constant stress amplitude of 430 MPa. Similar to AISI 304, it was observed that the nanocrystalline layer is stable up to 600 ◦ C under these conditions (Fig. 12). Moreover, the subsurface (5 m depth) shows a decrease of dislocation density with increasing temperature (Fig. 13). The microstructural changes in deep rolled AISI 304 as a function of number of cycles for a test temperature of 500 ◦ C and a stress amplitude of σ a = 280 MPa are shown in Figs. 14 and 15. Fracture in this case occurred after 4500
Fig. 10. Influence of test temperature on the half-width depth profile (σ a = 280 MPa, 2000 cycles).
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Fig. 11. TEM micrographs of the deep rolled and fatigued subsurface region of AISI 304 (5 m depth, σ a = 280 MPa, 2000 cycles, 100 ◦ C (left) and 600 ◦ C (right)).
cycles. The samples were fatigued and investigated after fixed loading cycles (1, 100, 1000 and 4000 cycles). Fig. 14 shows in an exemplary manner characteristic microstructures of the surface region and Fig. 15 of the subsurface region (5 m depth) for 1 and 4000 cycles corresponding to 0.02 and 88% of the total lifetime of the specimens. Again, the nanocrystalline layer is stable and the subsurface region is characterised by microstructural alterations and a decrease of dislocation density. At 4000 cycles (88% of fatigue lifetime) the subsurface showed tangled and dense dislocation arrangements. Diffuse cell structures parallel to the stress axis started to form (Fig. 15). In a previous study, the formation of a dislocation cell structure was observed in a deep rolled AISI 304 during cyclic deformation at high stress amplitudes [8]. It
can be summarised that isothermal fatigue up to 600 ◦ C (total time at elevated temperature of 17 min) does not affect the microstructure of the nanocrystalline layer, i.e. the shape and size of the grains both remained unchanged. The subsurface layer with its initially high dislocation density, however, is strongly affected and the dislocation density decreases during fatigue. As expected, the key parameter was the actual test temperature in these tests.
4. Discussion The fatigue life of metallic materials and components depends in a complex way on a multitude of factors among
Fig. 12. TEM micrographs of the deep rolled and fatigued surface region of Ti–6Al–4V (σ a = 430 MPa, 2000 cycles, 100 ◦ C (left) and 500 ◦ C (right)).
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Fig. 13. TEM micrograph of the deep rolled and fatigued subsurface of Ti–6Al–4V (σ a = 430 MPa, 2000 cycles, 100 ◦ C (left) and 500 ◦ C (right)).
which near-surface residual stresses and microstructures as well as roughness are just the most important ones for surface treated components. In assessing residual stress effects some considerable progress has been made over the past decades [19,22,25] and residual stress effects on fatigue performance have been thoroughly investigated even for complex parts and loading conditions [23,24,26]. Near-surface microstructures have been investigated much more rarely [9,15,19]. Moreover, surface treated components may experience alterations of near-surface properties during service load, especially under combined mechanical and thermal loading, such as in power plant or aircraft applications. The present study confirms that mechanical surface treatments, such as deep rolling, induce nanocrystalline
surfaces and highly strain hardened layers. The phenomenon of mechanically induced nanocrystallisation has also been observed in other contexts and related problems such as chip forming processes, hard turning or wear [29–31]. Nanocrystalline surface layers significantly affect the mechanical behaviour by restricting or impeding dislocation slip and the formation of slip bands at the surface [8], which act as preferential crack initiation sites. Therefore, taylored nanocrystalline surfaces are highly desirable in order to delay or prevent surface fatigue damage. One critical issue of mechanically induced nanocrystalline surface layers is their mechanical and thermal stability. Firstly, thermodynamically, nanocrystalline regions are in a non-equilibrium state with a high number of lattice defects that tend to anneal out, and finally recrystallise if
Fig. 14. TEM micrograph of the deep rolled and fatigued surface of AISI 304 (σ a = 280 MPa, 500 ◦ C, after own cycle (left) and 4000 cycles (right)).
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Fig. 15. TEM micrograph of the deep rolled and fatigued subsurface of AISI 304 (σ a = 280 MPa, 500 ◦ C, for 1 (left) and 4000 cycles (right)).
the temperatures are elevated enough to provide substantial diffusion processes. Secondly, according to recent in situ TEM studies [28], dislocation slip occurs in nanocrystalline materials, if the grain size is in the order of 20–50 nm or higher. Thus, it appeared therefore probable, that mechanical (cyclic or monotonic) loading affects the long-term or even short-term stability of nanocrystalline surface regions. By contrast, the present results indicate that nanocrystalline surface layers in AISI 304 and Ti–6Al–4V induced by deep rolling exhibit an extraordinarily high mechanical stability under the given conditions and remain very stable at high cyclic stress or plastic strain in the low-cycle fatigue (LCF) regime. This implies, that also in the high-cycle fatigue (HCF) regime where the plastic strain is much lower, mechanical stability of nanocrystalline surface layers is to be expected at low-test temperatures. The isothermal annealing experiments revealed that considerable grain coarsening takes place with increasing temperature and annealing time. In situ heating of TEM foils have yielded information on the thermal stability of nanocrystalline regions under shortterm annealing conditions [27]. An annealing temperature of 600 ◦ C can render the nanocrystalline surface layers “microcrystalline” after many hours of annealing as indicated by the present studies. Interestingly, also isothermal fatigue at elevated temperatures in the LCF regime (short LCF test) did not alter the structure of the nanocrystalline surface regions significantly in spite of simultaneous elevated temperature exposure and high cyclic plastic strain amplitude (which was in the order of 0.1–0.5%). Obviously, in the short LCF tests the exposure time is not sufficient to permit substantial annealing processes, therefore preventing coarsening and recrystallisation. We note that the results obtained from the presented investigations are not to be generalised for other test frequencies. The stability of nanocrystalline layers against high
temperature fatigue is clearly affected much more by the exposure time than by the number of cycles. Lower frequencies will inevitably lead to longer exposure time, and thus, to more pronounced diffusion processes and higher instability of the nanocrystallites. As a second consequence, instability of nanocrystalline layers may occur in the HCF regime, even though stability occurs in the LCF regime, because of different exposure times inspite of identical test temperatures. The more critical loading case may thus be the HCF case and not the LCF case. For high temperature fatigue it appears that the stability of nanocrystalline surface layers is mostly influenced by the temperature, followed by the exposure time, and finally by the stress/strain amplitude and number of cycles. Finally, the present study shows that the stability of nearsurface microstructures also strongly depends on the nature of near-surface microstructures: Underneath the nanocrystalline layer, the dislocation arrangements are generally much more instable to high temperature fatigue and transform into annealed low-energy arrangements with increasing temperature and increasing numbers of cycles as also indicated indirectly by FWHM value measurements.
5. Summary The mechanical, thermal and thermomechanical stability of nanostructured surface regions induced by deep rolling has been investigated in AISI 304 type metastable austenitic stainless steel and fan blade alloy Ti–6Al–4V. The results can be summarised as follows. • Deep rolling induced a nanocrystalline layer in both materials, deformation-induced martensite and mechanical twins in AISI 304 as well as high macro and micro residual stresses due to very high dislocation densities [19].
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• The nanocrystalline layer is stable during short time annealing up to 600 ◦ C. For AISI 304 long time annealing at 600 ◦ C (10–100 h) leads to recrystallisation. After 10 h at 600 ◦ C the nanocrystalline layer transforms into a fine equiaxed grain structure with an average grain size of about 150–200 nm. After 100 h at 600 ◦ C almost complete recrystallisation of the nanocrystalline layer takes place. • Isothermal fatigue in the LCF regime at high stress amplitudes does not significantly alter the structure of the nanocrystalline surface layer in the temperature range 25–600 ◦ C for the chosen test conditions. • Subsurface non-nanocrystalline regions with high dislocation densities tend to change into low energy dislocation arrangements during thermal, mechanical or thermomechanical loading. In all cases, a decline of the FWHM values of diffraction peaks during loading was observed.
Acknowledgement The authors would like to thank the German Science Foundation (DFG) for financial support of the Emmy-Noether group in Kassel (under contract-number AL 558/1-2).
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