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martensite multiphase steels
Materials Science & Engineering A 739 (2019) 404–414
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Effect of inclusion and microstructure on the very high cycle fatigue behaviors of high strength bainite/martensite multiphase steels
T
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Guhui Gaoa, , Qingzhen Xua, Haoran Guoa, Xiaolu Guia, Baoxiang Zhangb, Bingzhe Baic a
Material Science & Engineering Research Center, School of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing 100044, China Beijing General Research Institute of Nonferrous Metals, GRIKIN Adv Mat Co. Limited, Beijing 102200, China c Tsinghua University, Key Laboratory of Advanced Material, School of Material Science and Engineering, Beijing 100084, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Very high cycle fatigue Bainite Retained austenite Fracture mechanism Non-inclusion induced crack initiation
We studied the mutual and relative effect of inclusion and microstructure on very high cycle fatigue (VHCF) behavior of bainite/martensite (B/M) multiphase steel by designing four combinations of various inclusion sizes and microstructures. Results show that the fatigue crack initiations are mostly induced by inclusions in the specimens with large inclusion size, whereas the “non-inclusion induced crack initiations” (NIICI) are observed in the specimens whose microstructure is coarse. Unlike the inclusion induced cracks that grow along the plane of maximum tensile stress, the early non-inclusion induced cracks within micro-facet propagate along the plane of maximum shear stress (i.e., Stage I crack), which is related to the strain localization in oriented B/M laths during cyclic loading. Finally, with the assistance of scanning electron microscopy, electron backscatter diffraction, 3-Dimensional surface profile, focused ion beam and transmission electron microscopy, we proposed the NIICI mechanism of coarse B/M microstructure, and found the orientation and size of coarse B/M block and blocky retained austenite (RA) are crucial microstructural factors in facilitating the formation of NIICI within VHCF regime.
1. Introduction Very high cycle fatigue (VHCF, i.e., enduring cyclic stress in excess of 107 cycles) behavior of metallic materials has become an important subject to ensure the long-term safety of the structural components, such as aircraft (gas turbine disks 1010 cycles), automobiles (car engine 108 cycles), and railways (high speed train axles 109 cycles) [1–3]. Although a large amount of fatigue data has been published in the form of S-N curves, the data in the literature have been limited to fatigue lives up to 107 cycles. It is reported that fatigue failures of metallic material still happen beyond 107 cycles and the crack initiation site tends to transit from the specimen surface to the interior under a relatively low cyclic stress below the conventional fatigue limit [4–6]. Non-metallic inclusions acting as crack nucleation sites are usually observed in the center of a “bright granular facet” surrounded by a socalled “fish-eye” [2]. The bright granular facet (GBF) is also called optically dark areas, ODA [7] and fine granular area, FGA [1]. It is suggested that the non-metallic inclusions have a significant influence on the fatigue strength and fatigue life of high strength steels [8–10]. Hence, many works in VHCF behaviors of high-strength steels have focused on reducing the inclusions level and size by metallurgical
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methods to improve the VHCF properties [8]. Besides the inclusions, the microstructures also affect the VHCF behaviors of the metallic materials [11,12], in which case, the cracks initiates directly from microstructure rather than inclusions in high strength steels [13,14]. Previously, we defined this type of fatigue crack initiation as “non-inclusion induced crack initiation” (NIICI), compared with “inclusion induced crack initiation” (IICI) [13]. This phenomenon has been observed by some researchers. For instance, Chai first reported this phenomenon (termed as subsurface non-defect crack originsSNDFCO) in ferrite/martensite and austenite/martensite two-phase steels [15,16]. The formation of SNDFCO was caused by local plastic deformation of the soft phase (such as ferrite and austenite) [15]. Zhang et al. also found this phenomenon in very clean 54SiCr6 steel in which the inclusion size was smaller than 1 µm, but the crack initiation region was found to be enriched with carbon [17]. Our previous works indicated that the VHCF properties of the designed bainite/martensite (B/ M) multiphase steels were less sensitive to inclusions compared to tempered martensite steels. Hence, the NIICI became one of the most important fatigue failure modes in VHCF regime [18]. Yu et al. reported that formation of NIICI was related to the soft phase, such as ferrite and coarse bainitic ferrite lath, but direct evidence was still scant [18,19].
Correspondence to: School of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing, China. E-mail address: [email protected] (G. Gao).
https://doi.org/10.1016/j.msea.2018.10.073 Received 28 April 2018; Received in revised form 12 October 2018; Accepted 15 October 2018 Available online 16 October 2018 0921-5093/ © 2018 Elsevier B.V. All rights reserved.
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percentage (Fe balance). Manganese was added to avoid the austenite to ferrite transformation at high temperatures. Silicon was added to suppress carbide precipitation during quenching, partitioning and tempering steps. In the present work, two types of melting processes were employed: (i) melting by vacuum induction furnace; (ii) additional electroslag remelting after conventional vacuum melting, aiming to reduce the size of non-metallic inclusions in steel. For convenience, the steels without and with electroslag remelting were designated as V steel and E steel, respectively. All the steel ingots were reheated at 1200 °C and forged to 30 mm thickness with a finish-forging temperature at around 950 °C. The forged plate was then annealed at 950 °C followed by furnace cooling. The as-received microstructure prior to the following heat treatment was ferrite plus pearlite. The maximum inclusion size in steels after different melting processes was estimated by the statistics of the extreme values (SEV) method [14]. The inclusion size was observed directly from the inclusion induced fracture surface of V specimens, whereas about 30 standard inspection planes S0 with same size ~ 0.038 mm2 were chosen to observe inclusion size in E specimens. It is found that the cumulative probability for inclusion size can be well characterized by the Gumbel distribution, as shown in Fig. 1. Based on the Gumbel distribution, the evaluated maximum inclusion size in the smooth hour-glass type VHCF specimen is ~ 15 µm after electroslag remelting processes, which is much less than that (~ 98 µm) after conventional vacuum melting. The detailed calculation and data processing can be seen in Supplementary material. As shown by our previous work in Ref [20], two different BQ-P-T processes with different quenching and partitioning temperatures were carried out to tailor the microstructure: (i) cooling at ~ 2 °C/s to 200 °C (below Ms = 360 °C) after austenitization at 880 °C for 60 min, subsequent partitioning step at 280 °C for 45 min, and finally tempering at 250 °C for 120 min; (ii) cooling at ~ 2 °C/s to 320 °C (below Ms) after austenitization at 880 °C for 60 min, subsequent partitioning step at 360 °C for 45 min, and finally tempering at 250 °C for 120 min. The abbreviations of samples treated by various melting and BQ-P-T processes are summarized in Table 1. Microstructures and inclusions were characterized using optical microscopy, scanning electron microscope (SEM, ZEISS EVO18, 20 kV) with electron backscatter diffraction (EBSD, 20 kV, 200 nA, with step size: 0.1 µm and tilt angle: 70°) device and transmission electron microscopy (TEM, JEOL 2010, 200 kV) before and/or after fatigue test. The volume fraction of retained austenite (RA vol%) was measured by X-ray diffractometer (Rigaku Smartlab, Cu Kα radiation) at a step of 0.01° and a counting time of 2 s/step. The tensile properties and impact toughness of the experimental steel were measured at room temperature before fatigue test, and the
Recently, we investigated the crack initiation regions by EBSD analysis and found that the deformation-induced martensitic transformation (DIMT) of blocky retained austenite also induced the NIICI. The blocky retained austenite transforms easily to martensite due to local plastic deformation under cyclic loading, which is prone to induce the formation of micro-cracks [20]. But the origin of local plastic deformation and its correlation with microstructure are still unclear. P Zhao et al. observed the polished and etched fracture surface and found two interior non-inclusion crack initiation sites in the VHCF specimen with coarse microstructure [21]. However, since the original fracture surface was destroyed to some extent due to mechanical or electrolytic polishing, the detailed information about the nucleation and early growth of fatigue crack is missing. This brings about the challenge of understanding the microstructural nature of NIICI. The formation of GBF is generally considered as the early growth of fatigue crack, which is a dominant factor controlling the fatigue life. Until now, several mechanisms have been proposed to explain the formation of inclusion induced GBF, such as “hydrogen assisted crack growth” [22], “decohesion of spherical carbide” [23] and “formation and debonding of fine granular layer” [24]. All the mechanisms are based on the presence of non-metallic inclusions acting as pre-existing cracks or stress raisers. Since the pre-existing crack is absent, the formation of non-inclusion induced GBF involves both fatigue crack nucleation and early growth. This is also why we employ the term crack initiation rather than “nucleation” to refer to the processes. The microstructural characterization near non-inclusion induced GBF could provide an insight on the correlation between microstructure and crack initiation and help to understand the difference between the mechanisms of IICI and NIICI. In the present work, the effect of inclusions and microstructure on the VHCF behaviors (such as fatigue strength, life and crack initiation) was studied to explore the mechanisms of crack initiation from interior inclusions and microstructure. Here, we employed a novel heat treatment called quenching & partitioning (Q&P) process [25] to tailor the type, fraction, size and morphology of microstructure in steels with the same alloying elements. The microstructure features depend on the quenching temperature (Tq) and partitioning temperature (Tp) [26]. Moreover, the inclusion size was controlled by conventional vacuum melting and electroslag remelting processes (Seen in Fig. 1).
2. Experimental procedure 2.1. Test materials The chemical composition of the bainite/martensite (B/M) multiphase steel used was 0.2 C, 2.0Mn, 1.0Si, 0.8Cr, 0.2Mo, 0.6Ni in mass
Fig. 1. Gumbel distribution of inclusion size in (a) V specimens and (b) E specimens. 405
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Table 1 Mechanical, microstructural and inclusion properties of the four specimens treated by different melting and BQ-P-T processes. Samples
Smelting process
Heat treatment
Rm (MPa)
Rp (MPa)
A (%)
VQ200 EQ. 200 VQ320 EQ. 320
VM VM+ESR VM VM+ESR
Q200°C, Q200°C, Q320°C, Q320°C,
1402 1410 1328 1320
1080 ± 10 1130 ± 7 990 ± 5 995 ± 5
16.8 15.2 18.0 18.8
P280°C P280°C P360°C P360°C
± ± ± ±
15 3 12 10
± ± ± ±
0.3 0.5 0.5 0.3
AKV (J)
HV5
77 80 28 30
433 435 420 417
± ± ± ±
10 5 6 4
areaInc -max 117 µm 15 µm 117 µm 15 µm
Lblock 9.1 ± 3.8 µm 9.1 ± 3.8 µm 15.5 ± 3.6 µm 15.5 ± 3.6 µm
VM: Vacuum melting, ESR: electroslag remelting; Rm: tensile strength, Rp: yield strength, A: elongation, AKV: impact toughness, HV5: Vickers hardness; areaInc -max: maximum inclusion size, Lblock: length of B/M block.
ensures that the natural frequency of specimen matches the vibration frequency of ultrasonic testing machine. The VHCF testing was conducted on a Shimahzu USF-2000 type ultrasonic fatigue testing machine with a resonance frequency of 20 kHz and stress ratio of R = −1. Compressed air and intermittent loading (long interruption time: 600 ms) were introduced to relieve the possible “self-heating” in specimens during the VHCF testing [18]. However, it should be noted that the effect of heat generation cannot be impeded (only can be weakened) under ultrasonic tests at high stress amplitudes even by cooling the specimen and using a pulse-pause mode [28]. Hence, an “apparent VHCF strength” rather than “intrinsic VHCF strength” [18] was employed in this work. Crack initiation regions (e.g., fish eye and GBF areas) of fracture specimens were observed directly by SEM and 3-Dimensional Confocal Microscope Phase Shift MicroXAM-3D (Vertical scan resolution: 0.01 nm, lateral resolution: 0.11 µm). Here, MicroXAM-3D can measure the surface profile of the fatigue fracture surface. In order to investigate the microstructural evolution in fish eye and GBF areas, the fracture surface was etched directly in a 2 vol% Nital solution (named “etched” fracture surface), subsequent observed by SEM. Furthermore, the focused ion beam (FIB, Helios Nanolab 600i) equipment was carried out to investigate the relationship between microstructure and NIICI.
results are listed in Table 1. It is found that the melting processes have little influence on the conventional mechanical properties of the B/M steels. The mechanical properties, especially impact toughness, are closely related to the BQ-P-T process parameters, which decide the final microstructure of the B/M steels [27]. The final microstructures of the steels after different BQ-P-T treatment have been characterized by SEM, TEM and XRD in our previous work [20]. The smelting process has little influence on the microstructures (see in Supplementary material). Based on our previous work [20], the significant difference of microstructure between the two BQP-T treatments lays in the size, volume fraction and morphology of retained austenite (RA): nanometer-sized filmy RA in Q200 specimens but micrometer-sized blocky RA in Q320 counterparts. Here, the B/M blocks are also characterized by EBSD in this present work, as shown in Fig. 2. The EBSD results companied with pole figure analyses show that the B/M microstructures in both Q200 and Q320 specimens maintain well the Kurdjumov-Sachs (K-S) orientation relationship with their parent austenite [27]. However, almost all of 24 variants exist within a single prior austenite grain (PAG) in Q200 specimens, whereas there are only 4 or 6 variants within an individual PAG in Q320 specimens. It means that more numbers and consequent smaller size of block are formed in Q200 specimens [27]. Through the interception line method, the measured average length of block in Q200 specimens is 9.1 ± 3.8 µm, which is much less than that (15.5 ± 3.6 µm) in Q320 counterparts. In summary, four specimens with various microstructure and inclusion sizes were designed in the present work: VQ200 specimen with large inclusion size but fine microstructure (refined B/M block and nano-sized filmy RA); EQ. 200 specimen with small inclusion size and fine microstructure; VQ320 specimen with large inclusion size and coarse microstructure; EQ. 320 specimen with small inclusion size but coarse microstructure. The mechanical, microstructural and inclusion properties of these four specimens are listed in Table 1.
3. Results 3.1. S-N characteristics Fig. 3a-d shows the S-N data obtained from VHCF testing of the four B/M steel specimens. The apparent mean fatigue strength [18] was determined by the conventional staircase method (i.e., up and down method, seen in supplementary material) by setting the run-out number of cycle near 109 [29]. The abbreviations in the upper right corner mean the crack initiation modes. The apparent mean fatigue strength (σw9) of the VQ200, EQ. 200, VQ320 and EQ. 320 specimens are 592 ± 47 MPa, 768 ± 13 MPa, 570 ± 44 MPa and 603 ± 13 MPa, respectively. The ratios of apparent mean fatigue strength to tensile strength (σw9/Rm) of the four variants are 0.42 ± 0.03, 0.52 ± 0.01, 0.43 ± 0.03 and 0.45 ± 0.01, respectively (Fig. 3f). It can be found that the inclusion size and microstructure have a
2.2. Fatigue testing method The geometry and dimensions of the smooth hour-glass type specimens for VHCF testing are shown in Ref [20]. The specimen geometry is dependent on the Young's modulus and density of steel, which
Fig. 2. EBSD observation showing the size and orientation of bainite/martensite blocks in (a) EQ. 200 and (b) EQ. 320 samples; bold black lines: misorientation, θ ≥ 15°, white lines: prior austenite grain boundaries, the insets showing {0 0 1} pole figure. 406
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Fig. 3. S-N data of the four B/M multiphase steels treated by various melting and BQ-P-T processes, (a) VQ200, (b) EQ. 200, (c) VQ320, (d) EQ. 320, and (e) the change of fatigue strength as well as (f) the ratio ofσw9/Rm for the four specimens; Sur: surface induced crack initiation, Sur (Inc): surface inclusion induced crack initiation, Int (Inc): interior inclusion induced crack initiation, Int (Non-inc): interior non-inclusion induced crack initiation (crack initiated from microstructure) [20].
3.2. Fracture surface observation
mutual and relative effect on the VHCF property of B/M steel, in other word, the effect of inclusion size on the fatigue strength strongly depends on the microstructure, and vice versa. For instance, the reduction of inclusion size can bring about 176 MPa increment of σw9 for the specimens with fine microstructure (VQ200 and EQ. 200), whereas the increment is only 33 MPa for the specimens with coarse microstructure (VQ320 and EQ. 320), as indicated by red arrows in Fig. 3e. Similarly, the refinement of microstructure has little influence on fatigue strength when the inclusion size is large, but the advantage of microstructural refinement becomes exciting when the inclusion size is reduced, as indicated by black arrows in Fig. 3e.
Based on fracture surface observation, the fatigue crack initiations are mostly induced by inclusions in the specimens with large inclusion size, e.g., 100% (11/11) in VQ200 specimen and 86.7% (13/15) in VQ320 variant. In contrast, no IICI is found in the specimens with small inclusion size due to electroslag remelting. But the fatigue cracks could initiate directly from interior microstructure when the microstructure is coarse, e.g., 13.3% (2/15) in VQ320 specimen and 92.3% (12/13) in EQ. 320 variant. Fig. 4a and b show the crack initiations are induced by surface inclusions or interior inclusions in VQ200 specimen. The GBF can be observed around the interior inclusions (Fig. 4b), but it is not evident 407
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Fig. 4. SEM images showing fractography of the VQ200 specimen, (a) surface inclusion (σa = 650 MPa, Nf = 8.75 × 105 cycles), (b) interior inclusion with GBF (σa = 575 MPa, Nf = 8.99 × 107 cycles).
Fig. 5. SEM images showing a “Non-inclusion induced crack initiation” fracture surface of the VQ320 specimen (σa = 600 MPa, Nf = 9.44 × 106 cycles) and (b) is the detailed observation of GBF; the cross in (b) showing the EDS gathering area, the inset in (b) showing the EDS spectrum.
Fig. 6. SEM images showing a “Non-inclusion induced crack initiation” fracture surface of the EQ. 320 specimen (σa = 650 MPa, Nf = 9.51 × 106 cycles) and (b) is the detailed observation of GBF.
which will be discussed in the following. Besides, only surface induced fracture is observed and no fatigue failure occurs at number of cycles exceeding 107 for EQ. 200 specimen [20], as shown in Fig. 3b. U. Krupp et al. [30] also reported the similar phenomenon and proposed that the absence of the failure at number of cycles exceeding 107 indicated that the early micro-cracks are nonpropagating cracks and permanently blocked by microstructural barriers. It is implied that, different from the other three specimens, neither interior inclusions nor microstructure could induce the crack initiation during VHCF regime for EQ. 200 specimen.
around the surface inclusions (Fig. 4a). Fig. 5 and Fig. 6 show the fracture surfaces of NIICI in the VQ320 and EQ. 320 samples, respectively. The fish-eye with GBF is also formed in the non-inclusion induced fracture surface. However, the detailed EDS analysis indicates that no inclusions exist around the GBF and the composition in the noninclusion induced GBF is similar to that in the matrix, which confirms that the crack initiation is induced by interior microstructure rather than inclusions [20]. The SEM observation in Fig. 6b shows the presence of numerous submicron-sized facets and the trace of severe plastic deformation within GBF. The severe plastic deformation might come from the rubbing of two fracture surface or the strain localization, or both. The former is generated after the formation of cracks, whereas the later occurs before crack initiation and may be the cause of NIICI, 408
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3.3. Stress intensity factor range The parameters of characteristic area can be obtained from fracture surface observation. Then, the value of stress intensity factor (SIF) range at the periphery of inclusion (ΔKInc), GBF around inclusion (ΔKIG) and GBF without inclusion (ΔKNG) are calculated based on the area model [31]:
∆K = 0.5σa π area
(1)
where σa is stress amplitude, area is the projected area of inclusion and GBF on the plane perpendicular to tensile stress and area is used to characterized the size of inclusion and GBF [32]. For surface inclusion or defect, the coefficient of 0.5 is replaced by 0.65 in the calculation [32]. Fig. 7 shows the calculated value of ΔKInc, ΔKIG and ΔKNG versus fatigue life. The values of ΔKInc are between 2.0 and 4.5 MPa√m, with the value of areaInc ranging from 10 µm to 83 µm (Fig. 1a). The values of ΔKIG are between 3.3 and 6.6 MPa√m, while those of ΔKNG are relatively small, varying from 2.9 to 5.0 MPa√m. The important point was that the values of ΔKIG and ΔKNG almost keep constant at average of 4.8 MPa√m and 3.8 MPa√m under the failure cycles between 106 and 109, whereas the values of ΔKInc display a slightly decreasing tendency with respect to failure life. This confirms the correlation between inclusion size and fatigue life.
Fig. 7. Stress intensity factor range of inclusion (Inc), GBF with inclusion (IGBF) and GBF without inclusion (N-GBF) of the VQ200, VQ320 and EQ. 320 specimens.
Fig. 8. MicroXAM-3D observation of an “inclusions induced crack initiation” fracture surface in VQ320 specimen (σa = 600 MPa, Nf = 1.47 × 108 cycles), (a) the optical image, the inset in (a) showing the EDS spectrum, (b) 3D morphology image, (c) the sectional profile of the labeled line in (b), the bottom is an enlarged window showing the detailed observation of GBF. 409
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Fig. 9. MicroXAM-3D observation of an “Non-inclusions induced crack initiation” fracture surface in EQ. 320 specimen (σa = 650 MPa, Nf = 9.51 × 106 cycles), (a) the optical image, (b) 3D morphology image, (c) the sectional profile of the labeled line in (b), the bottom is an enlarged window showing the detailed observation of GBF. (Corresponding SEM observation is shown in Fig. 6).
strength strongly depends on the microstructure, and vice versa. For the refined microstructure, the reduction of inclusion size could bring an exciting increment of fatigue strength, but the efficacy is weak for the coarse microstructure. It is because, although the IICI is eliminated due to the reduction of inclusion size, the NIICI would become the dominant fracture mode when the microstructure is coarse (e.g., in EQ. 320 specimen). In other word, there is a competition process between nonmetallic inclusion and microstructure to determining the VHCF properties of B/M steels [34]. Hence, it is essential to understand the mechanisms of IICI and NIICI to explore the crucial factors in controlling the fracture modes. Since the mechanism of IICI has been reported in many literatures [22–24], we will focus on the mechanism of NIICI in the following. Through the surface profiles analysis by MicroXAM-3D, we find that, at macroscopic scale, the surface of inclusion induced GBF (IG) is normal to the loading direction (Fig. 8), but the non-inclusion induced GBF (NG) is along close to the plane of maximum shear stress (Figs. 9 and 10). The surface profiles of NG have a strong similarity with the stage I crack propagation, which is identified in single crystals with crack growing in a plane oriented at an angle about 45° against the loading direction [35]. Stage I crack propagation is also typical of the early propagation of micro-cracks in surface grains as shown by Forsyth [36]. This stage is sensitive to microstructure and related to slip/strain localization in individual grains [35]. Owing to the irreversible dislocation slip, the persistent slip bands (PSBs) accompanied with
3.4. Fracture mechanisms There are two dominant interior crack initiation modes for B/M steels within VHCF regimes, namely, IICI and NIICI modes. In order to investigate the difference between the fracture mechanisms of two crack initiation modes, the MicroXAM-3D was employed to measure the surface profile of the fatigue fracture surface. Fig. 8 shows the IICI fracture surface profile in the VQ320 specimen. The failure is caused by interior inclusion, and the GBF with fish eye can be observed around the inclusion (the GBF is also called optically dark area, ODA, Fig. 8a). As shown by the profile (Fig. 8c), the surfaces of GBF and fisheye are all perpendicular to the loading direction, which indicates that the early cracks in GBF are formed directly due to inclusions and then grow along the plane of maximum tensile stress [33]. Figs. 9 and 10 show the NIICI fracture surface profiles of the EQ. 320 specimen. The surface profiles show the GBF actually contains a micro-facet with size of tens of micrometer. This facet has an angle of about 40° against the plane of maximum tensile stress, namely, close to the plane of maximum shear stress. It could be inferred that, unlike inclusion induced cracks, the early propagation of the non-inclusion induced cracks within GBF is controlled by maximum shear stress [33]. 4. Discussions As indicated in Fig. 3, the effect of inclusion size on the fatigue 410
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Fig. 10. MicroXAM-3D observation of an “Non-inclusions induced crack initiation” fracture surface in EQ. 320 specimen (σa = 675 MPa, Nf = 1.89 × 106 cycles), (a) the optical image, (b) 3D morphology image, (c) the sectional profile of the labeled line in (b), the bottom is an enlarged window showing the detailed observation of GBF.
micro-facet through continued cyclic loading. Another question arises: since the orientation of bainite/martensite lath is on the whole random (in whether refined or coarse microstructure), the conditions of strain localization could be fulfilled for any lath along 45° oriented {110} plane against the loading direction [38]. However, we observe the NIICI only in the Q320 specimens with coarse microstructure. In order to explain this phenomenon, we observe the NIICI fracture surface on a relatively large scale. Fig. 12 shows a typical NIICI fracture surfaces before and after etching of an EQ. 320 specimen failed at Nf = 7.62 × 107 with the stress amplitude σa = 600 MPa. It can be found that the region of fish eye envelops a GBF (equivalent length 46.2 µm), but no inclusion is found around GBF (Fig. 12a and b). After etching, the micro-facet is obviously observed, whose size is smaller than that of GBF (Fig. 12c). Through backscattered electron image (Fig. 12d), we find that the micro-facet basically runs through the prior austenite grain (PAG) and directly impinges on the PAG boundaries. And the micro-voids and micro-cracks are observed along the meeting place of micro-facet and PAG boundaries. It implies that the orientation and dimension of microfacet and the state of PAG boundaries are crucial factors in determining whether the strain localization could develop the NIICI crack. Based on the discussion above, the mechanism of NIICI in VHCF regime for the coarse B/M microstructure is suggested:
extrusion and intrusions lead to the surface relief which acts as stress raiser due to its notch effect in the continuation of the cyclic deformation and as a consequence can lead to crack initiation [37]. In the present work, it is interesting that the Stage I crack propagation is found in the interior microstructure, but in which case, the relief at free surface cannot form. Then, a question arises that how the crack initiates from interior microstructure. We carried out the focused ion beam (FIB) sample cut from the micro-facet to observe the correlation of microstructural characteristic and crack initiation origin, as shown in Fig. 11. It can be found that the profile of FIB sample (Fig. 11c) is consistent with the results of MicroXAM-3D (Figs. 7c and 8c), which confirms that the micro-facet is along the plane of maximum shear stress. A sheaf of parallel bainite/ martensite lath is observed beneath the micro-facet, and the micro-facet lies parallel to the lath longitudinal direction. Furthermore, the selected area diffraction (SAD) patterns show that the {110} planes of the B/M lath is also along the direction of maximum shear stress (Fig. 11d). L. Morsdorf et al. [38] studied the plasticity of lath martensite by EBSD, digital image correlation (DIC) and electron channeling contrast imaging (ECCI) and found that the slip in the lath preferred to activate along the 45° oriented {110} plane against the loading direction, where the shear stress is highest. Hence, the localized slip/strain could firstly occur along ~ 45° oriented laths, which could result in the formation of 411
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Fig. 11. Focused ion beam (FIB) observation showing the correlation of microstructural characteristic and crack initiation of VQ320 sample (σa = 600 MPa, Nf = 4.38 × 107), (a) SEM image showing the GBF area, (b) SEM image showing the location of FIB cut, Pt coating was to protect the fracture surface, (c) SEM image showing the profile of fracture surface, (d) TEM image and SAD showing the microstructure beneath GBF and the orientation of B/M lath.
Fig. 12. SEM image showing typical “non-inclusion induced crack initiation” fracture surfaces before (a, b) and after etching (c, d) of an EQ. 320 specimen (σa = 600 MPa, Nf = 7.62 × 107), PAGB: prior austenite grain boundary, RA: retained austenite. 412
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Fig. 13. The schematic representation of NIICI mechanism in coarse B/M microstructure, the occurrence of NIICI progresses according to the order of (a) → (b) → (c) → (d), PAGB: prior austenite grain boundary, RA: retained austenite, M: newly formed martensite from RA under cyclic loading.
Furthermore, the nano-sized filmy RA located between B/M lathes has high stability and also acts as strong barrier to slip movement and crack propagation due to the dislocation absorption and transformation induced plasticity effects [45,46]. In this case, the early localized slip bands are permanently blocked by these microstructural barriers, and consequently the NIICI cannot occur in the Q200 specimen with refined microstructure.
The first stage (Fig. 13a): under cyclic loading, slip localization prefers to occur along some certain oriented crystallographic planes [39], e.g., the {110} planes oriented in 45° against the loading direction (as shown in Fig. 11d). The bainite and martensite in low carbon steels have a three-level hierarchy in its morphology, i.e. packet, block and lath [40]. The packet is a group of laths with a common habit plane with respect to austenite. Each packet is divided into blocks, which contain the parallel laths with the same crystallographic orientation and is also called as a co-variant packet. Since the micro-facet is parallel to the longitudinal direction of the B/M laths with the same {110} plane oriented in 45° against the loading direction (Fig. 11c), it is deduced that the length of blocks governs the slip localization at the early stage of fatigue for B/M steels [41]. This was also confirmed by Mueller et al. [42] through building the constitutive relationship between block size and fatigue limit in nano-structured bainitic steels. The second stage (Fig. 13b): with increasing cycles, in case that the B/M block is coarse and grows through the PAG (e.g., the B/M block in Q320 specimen, Fig. 2b), the localized slip bands could directly impinge on the packet or PAG boundaries without any microstructural barrier. At the third stage (Fig. 13c), whether or not the localized slip bands could grow across the PAG boundaries into the adjacent PAG depends on the state of PAG boundaries and the critical cyclic stress for dislocation moving [43]. Since most of blocky RA is located nearby the PAG boundaries in EQ. 320 specimen [20], the strength of PAG would be reduced. Furthermore, the blocky RA along PAG boundaries transforms easily to brittle martensite due to localized plastic deformation under the impinging of slip bands on PAG boundaries [20]. The brittle martensite is prone to induce the formation of cleavage micro-cracks along the PAG boundaries (as shown in Fig. 12d). At the final stage (Fig. 13d), the micro-cracks and voids along the PAG boundaries act as stress raisers and help the localized slip bands grow across the PAG boundaries, and then the occurrence of NIICI becomes possible. The schematic representation of NIICI mechanism in coarse B/M microstructure is shown in Fig. 13. In contrast, the orientation distribution of refined B/M blocks are intersecting in Q200 sample (as shown in Fig. 2a), which supplies more boundary barriers to restricting the localized slip progression [44].
5. Conclusions In this work, four B/M steel specimens with various microstructure and inclusion size were prepared by controlling the melting and BQ-P-T processes, i.e., VQ200 specimen with large inclusion size (maximum size: ~ 98 µm) but fine microstructure (refined B/M block and nanosized filmy RA); EQ. 200 specimen with small inclusion size (maximum size: ~ 15 µm) and fine microstructure; VQ320 specimen with large inclusion size and coarse microstructure; EQ. 320 specimen with small inclusion size but coarse microstructure. The effect of inclusions and microstructure on the very high cycle fatigue behaviors of B/M multiphase steel was investigated through S-N curves testing, fracture surface observation, failure mechanism and crack initiation analyses. In particular, the mechanism of “non-inclusion induced crack initiations” was investigated. The main conclusions are described as follows: (1) The inclusion size and microstructure have a mutual and relative effect on the VHCF property of B/M steel, i.e., the effect of inclusion size on the fatigue strength strongly depends on the microstructure, and vice versa. (2) The “inclusion induced crack initiation” (IICI) is a leading failure mode within VHCF regimes for the specimens with large inclusion size, while we observe the “non-inclusion induced crack initiations” (NIICI) only in the specimens whose microstructure is coarse. (3) At macroscopic scale, the surface of inclusion induced GBF is normal to the loading direction, but the non-inclusion induced GBF is along close to the plane of maximum shear stress. (4) The NIICI micro-facet is parallel to the longitudinal direction of the B/M laths with {110} plane oriented in 45° against the loading 413
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direction，which is related to the strain localization during cyclic loading. And the micro-voids and micro-cracks are observed along the meeting place of NIICI micro-facet and PAG boundaries. (5) The orientation and size of coarse B/M block and blocky RA are crucial microstructural factors in facilitating the formation of NIICI.
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Acknowledgment This research is supported by National Natural Science Foundation of China (No. 51771014 and U1760104), the Fundamental Research Funds for the Central Universities (No.2017JBM053), and Natural Science Foundation of Beijing (2172047). The authors would also like to express their appreciation to Central Iron & Steel Research Institute for supporting this study and providing ultrasonic fatigue testing equipment (SHIMADZU USF-2000). Declarations of interest None. Appendix A. Supporting information Supplementary data associated with this article can be found in the online version at doi:10.1016/j.msea.2018.10.073. References [1] T. Sakai, Y. Sato, N. Oguma, Characteristic S–N properties of high‐carbon–chromium‐bearing steel under axial loading in long‐life fatigue, Fatigue Fract. Eng. Mater. 25 (8‐9) (2002) 765–773. [2] K. Shiozawa, L. Lu, Very high‐cycle fatigue behaviour of shot‐peened high‐carbon–chromium bearing steel, Fatigue Fract. Eng. Mater. 25 (8‐9) (2002) 813–822. [3] C. Bathias, C.P. Paris, Gigacycle Fatigue in Mechanical Practice, Taylor & Francis, 2004. [4] B. Pyttel, D. Schwerdt, C. Berger, Very high cycle fatigue–is there a fatigue limit? Int. J. Fatigue 33 (1) (2011) 49–58. [5] Q. Jiang, C. Sun, X. Liu, Y. Hong, Very-high-cycle fatigue behavior of a structural steel with and without induced surface defects, Int. J. Fatigue 93 (2016) 352–362. [6] T. Sakai, A. Nakagawa, N. Oguma, Y. Nakamura, A. Ueno, S. Kikuchi, A. Sakaida, A review on fatigue fracture modes of structural metallic materials in very high cycle regime, Int. J. Fatigue 93 (2016) 339–351. [7] K. Shiozawa, Y. Morii, S. Nishino, L. Lu, Subsurface crack initiation and propagation mechanism in high-strength steel in a very high cycle fatigue regime, Int. J. Fatigue 28 (2006) 1521–1532. [8] S.X. Li, Effects of inclusions on very high cycle fatigue properties of high strength steels, Int Mater. Rev. 57 (2) (2012) 92–114. [9] A. Pineau, S. Forest, Effects of inclusions on the very high cycle fatigue behaviour of steels, Fatigue Fract. Eng. Mater. 40 (11) (2017) 1694–1707. [10] U. Karr, R. Schuller, M. Fitzka, B. Schönbauer, D. Tran, B. Pennings, H. Mayer, Influence of inclusion type on the very high cycle fatigue properties of 18Ni maraging steel, J. Mater. Sci. 52 (10) (2017) 5954–5967. [11] D.L. McDowell, F.P.E. Dunne, Microstructure-sensitive computational modeling of fatigue crack formation, Int J. Fatigue 32 (9) (2010) 1521–1542. [12] G.M. Castelluccio, W.D. Musinski, D.L. McDowell, Recent developments in assessing microstructure-sensitive early stage fatigue of polycrystals, Curr. Opin. Solid State Mater. 18 (4) (2014) 180–187. [13] P. Zhao, G. Gao, R.D.K. Misra, B. Bai, Effect of microstructure on the very high cycle fatigue behavior of a bainite/martensite multiphase steel, Mater. Sci. Eng. A Struct. 630 (2015) 1–7. [14] W. Li, T. Sakai, M. Wakita, S. Mimura, Influence of microstructure and surface defect on very high cycle fatigue properties of clean spring steel, Int. J. Fatigue 60 (2014) 48–56. [15] G. Chai, The formation of subsurface non-defect fatigue crack origins, Int. J. Fatigue 28 (11) (2006) 1533–1539. [16] G. Chai, N. Zhou, Study of crack initiation or damage in very high cycle fatigue using ultrasonic fatigue test and microstructure analysis, Ultrasonics 53 (8) (2013) 1406–1411. [17] J. Zhang, S. Li, Z. Yang, Z. Yang, G. Li, W. Hui, Y. Weng, Influence of inclusion size on fatigue behavior of high strength steels in the gigacycle fatigue regime, Int. J. Fatigue 29 (4) (2007) 765–771. [18] Y. Yu, J. Gu, L. Xu, F. Shou, B. Bai, Y. Liu, Very high cycle fatigue behaviors of Mn–Si–Cr series bainite/martensite dual phase steels, Mater. Des. 31 (6) (2010)
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Effect of inclusion and microstructure on the very high cycle fatigue behaviors of high strength bainite/martensite multiphase steels
T
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Guhui Gaoa, , Qingzhen Xua, Haoran Guoa, Xiaolu Guia, Baoxiang Zhangb, Bingzhe Baic a
Material Science & Engineering Research Center, School of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing 100044, China Beijing General Research Institute of Nonferrous Metals, GRIKIN Adv Mat Co. Limited, Beijing 102200, China c Tsinghua University, Key Laboratory of Advanced Material, School of Material Science and Engineering, Beijing 100084, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Very high cycle fatigue Bainite Retained austenite Fracture mechanism Non-inclusion induced crack initiation
We studied the mutual and relative effect of inclusion and microstructure on very high cycle fatigue (VHCF) behavior of bainite/martensite (B/M) multiphase steel by designing four combinations of various inclusion sizes and microstructures. Results show that the fatigue crack initiations are mostly induced by inclusions in the specimens with large inclusion size, whereas the “non-inclusion induced crack initiations” (NIICI) are observed in the specimens whose microstructure is coarse. Unlike the inclusion induced cracks that grow along the plane of maximum tensile stress, the early non-inclusion induced cracks within micro-facet propagate along the plane of maximum shear stress (i.e., Stage I crack), which is related to the strain localization in oriented B/M laths during cyclic loading. Finally, with the assistance of scanning electron microscopy, electron backscatter diffraction, 3-Dimensional surface profile, focused ion beam and transmission electron microscopy, we proposed the NIICI mechanism of coarse B/M microstructure, and found the orientation and size of coarse B/M block and blocky retained austenite (RA) are crucial microstructural factors in facilitating the formation of NIICI within VHCF regime.
1. Introduction Very high cycle fatigue (VHCF, i.e., enduring cyclic stress in excess of 107 cycles) behavior of metallic materials has become an important subject to ensure the long-term safety of the structural components, such as aircraft (gas turbine disks 1010 cycles), automobiles (car engine 108 cycles), and railways (high speed train axles 109 cycles) [1–3]. Although a large amount of fatigue data has been published in the form of S-N curves, the data in the literature have been limited to fatigue lives up to 107 cycles. It is reported that fatigue failures of metallic material still happen beyond 107 cycles and the crack initiation site tends to transit from the specimen surface to the interior under a relatively low cyclic stress below the conventional fatigue limit [4–6]. Non-metallic inclusions acting as crack nucleation sites are usually observed in the center of a “bright granular facet” surrounded by a socalled “fish-eye” [2]. The bright granular facet (GBF) is also called optically dark areas, ODA [7] and fine granular area, FGA [1]. It is suggested that the non-metallic inclusions have a significant influence on the fatigue strength and fatigue life of high strength steels [8–10]. Hence, many works in VHCF behaviors of high-strength steels have focused on reducing the inclusions level and size by metallurgical
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methods to improve the VHCF properties [8]. Besides the inclusions, the microstructures also affect the VHCF behaviors of the metallic materials [11,12], in which case, the cracks initiates directly from microstructure rather than inclusions in high strength steels [13,14]. Previously, we defined this type of fatigue crack initiation as “non-inclusion induced crack initiation” (NIICI), compared with “inclusion induced crack initiation” (IICI) [13]. This phenomenon has been observed by some researchers. For instance, Chai first reported this phenomenon (termed as subsurface non-defect crack originsSNDFCO) in ferrite/martensite and austenite/martensite two-phase steels [15,16]. The formation of SNDFCO was caused by local plastic deformation of the soft phase (such as ferrite and austenite) [15]. Zhang et al. also found this phenomenon in very clean 54SiCr6 steel in which the inclusion size was smaller than 1 µm, but the crack initiation region was found to be enriched with carbon [17]. Our previous works indicated that the VHCF properties of the designed bainite/martensite (B/ M) multiphase steels were less sensitive to inclusions compared to tempered martensite steels. Hence, the NIICI became one of the most important fatigue failure modes in VHCF regime [18]. Yu et al. reported that formation of NIICI was related to the soft phase, such as ferrite and coarse bainitic ferrite lath, but direct evidence was still scant [18,19].
Correspondence to: School of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing, China. E-mail address: [email protected] (G. Gao).
https://doi.org/10.1016/j.msea.2018.10.073 Received 28 April 2018; Received in revised form 12 October 2018; Accepted 15 October 2018 Available online 16 October 2018 0921-5093/ © 2018 Elsevier B.V. All rights reserved.
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percentage (Fe balance). Manganese was added to avoid the austenite to ferrite transformation at high temperatures. Silicon was added to suppress carbide precipitation during quenching, partitioning and tempering steps. In the present work, two types of melting processes were employed: (i) melting by vacuum induction furnace; (ii) additional electroslag remelting after conventional vacuum melting, aiming to reduce the size of non-metallic inclusions in steel. For convenience, the steels without and with electroslag remelting were designated as V steel and E steel, respectively. All the steel ingots were reheated at 1200 °C and forged to 30 mm thickness with a finish-forging temperature at around 950 °C. The forged plate was then annealed at 950 °C followed by furnace cooling. The as-received microstructure prior to the following heat treatment was ferrite plus pearlite. The maximum inclusion size in steels after different melting processes was estimated by the statistics of the extreme values (SEV) method [14]. The inclusion size was observed directly from the inclusion induced fracture surface of V specimens, whereas about 30 standard inspection planes S0 with same size ~ 0.038 mm2 were chosen to observe inclusion size in E specimens. It is found that the cumulative probability for inclusion size can be well characterized by the Gumbel distribution, as shown in Fig. 1. Based on the Gumbel distribution, the evaluated maximum inclusion size in the smooth hour-glass type VHCF specimen is ~ 15 µm after electroslag remelting processes, which is much less than that (~ 98 µm) after conventional vacuum melting. The detailed calculation and data processing can be seen in Supplementary material. As shown by our previous work in Ref [20], two different BQ-P-T processes with different quenching and partitioning temperatures were carried out to tailor the microstructure: (i) cooling at ~ 2 °C/s to 200 °C (below Ms = 360 °C) after austenitization at 880 °C for 60 min, subsequent partitioning step at 280 °C for 45 min, and finally tempering at 250 °C for 120 min; (ii) cooling at ~ 2 °C/s to 320 °C (below Ms) after austenitization at 880 °C for 60 min, subsequent partitioning step at 360 °C for 45 min, and finally tempering at 250 °C for 120 min. The abbreviations of samples treated by various melting and BQ-P-T processes are summarized in Table 1. Microstructures and inclusions were characterized using optical microscopy, scanning electron microscope (SEM, ZEISS EVO18, 20 kV) with electron backscatter diffraction (EBSD, 20 kV, 200 nA, with step size: 0.1 µm and tilt angle: 70°) device and transmission electron microscopy (TEM, JEOL 2010, 200 kV) before and/or after fatigue test. The volume fraction of retained austenite (RA vol%) was measured by X-ray diffractometer (Rigaku Smartlab, Cu Kα radiation) at a step of 0.01° and a counting time of 2 s/step. The tensile properties and impact toughness of the experimental steel were measured at room temperature before fatigue test, and the
Recently, we investigated the crack initiation regions by EBSD analysis and found that the deformation-induced martensitic transformation (DIMT) of blocky retained austenite also induced the NIICI. The blocky retained austenite transforms easily to martensite due to local plastic deformation under cyclic loading, which is prone to induce the formation of micro-cracks [20]. But the origin of local plastic deformation and its correlation with microstructure are still unclear. P Zhao et al. observed the polished and etched fracture surface and found two interior non-inclusion crack initiation sites in the VHCF specimen with coarse microstructure [21]. However, since the original fracture surface was destroyed to some extent due to mechanical or electrolytic polishing, the detailed information about the nucleation and early growth of fatigue crack is missing. This brings about the challenge of understanding the microstructural nature of NIICI. The formation of GBF is generally considered as the early growth of fatigue crack, which is a dominant factor controlling the fatigue life. Until now, several mechanisms have been proposed to explain the formation of inclusion induced GBF, such as “hydrogen assisted crack growth” [22], “decohesion of spherical carbide” [23] and “formation and debonding of fine granular layer” [24]. All the mechanisms are based on the presence of non-metallic inclusions acting as pre-existing cracks or stress raisers. Since the pre-existing crack is absent, the formation of non-inclusion induced GBF involves both fatigue crack nucleation and early growth. This is also why we employ the term crack initiation rather than “nucleation” to refer to the processes. The microstructural characterization near non-inclusion induced GBF could provide an insight on the correlation between microstructure and crack initiation and help to understand the difference between the mechanisms of IICI and NIICI. In the present work, the effect of inclusions and microstructure on the VHCF behaviors (such as fatigue strength, life and crack initiation) was studied to explore the mechanisms of crack initiation from interior inclusions and microstructure. Here, we employed a novel heat treatment called quenching & partitioning (Q&P) process [25] to tailor the type, fraction, size and morphology of microstructure in steels with the same alloying elements. The microstructure features depend on the quenching temperature (Tq) and partitioning temperature (Tp) [26]. Moreover, the inclusion size was controlled by conventional vacuum melting and electroslag remelting processes (Seen in Fig. 1).
2. Experimental procedure 2.1. Test materials The chemical composition of the bainite/martensite (B/M) multiphase steel used was 0.2 C, 2.0Mn, 1.0Si, 0.8Cr, 0.2Mo, 0.6Ni in mass
Fig. 1. Gumbel distribution of inclusion size in (a) V specimens and (b) E specimens. 405
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Table 1 Mechanical, microstructural and inclusion properties of the four specimens treated by different melting and BQ-P-T processes. Samples
Smelting process
Heat treatment
Rm (MPa)
Rp (MPa)
A (%)
VQ200 EQ. 200 VQ320 EQ. 320
VM VM+ESR VM VM+ESR
Q200°C, Q200°C, Q320°C, Q320°C,
1402 1410 1328 1320
1080 ± 10 1130 ± 7 990 ± 5 995 ± 5
16.8 15.2 18.0 18.8
P280°C P280°C P360°C P360°C
± ± ± ±
15 3 12 10
± ± ± ±
0.3 0.5 0.5 0.3
AKV (J)
HV5
77 80 28 30
433 435 420 417
± ± ± ±
10 5 6 4
areaInc -max 117 µm 15 µm 117 µm 15 µm
Lblock 9.1 ± 3.8 µm 9.1 ± 3.8 µm 15.5 ± 3.6 µm 15.5 ± 3.6 µm
VM: Vacuum melting, ESR: electroslag remelting; Rm: tensile strength, Rp: yield strength, A: elongation, AKV: impact toughness, HV5: Vickers hardness; areaInc -max: maximum inclusion size, Lblock: length of B/M block.
ensures that the natural frequency of specimen matches the vibration frequency of ultrasonic testing machine. The VHCF testing was conducted on a Shimahzu USF-2000 type ultrasonic fatigue testing machine with a resonance frequency of 20 kHz and stress ratio of R = −1. Compressed air and intermittent loading (long interruption time: 600 ms) were introduced to relieve the possible “self-heating” in specimens during the VHCF testing [18]. However, it should be noted that the effect of heat generation cannot be impeded (only can be weakened) under ultrasonic tests at high stress amplitudes even by cooling the specimen and using a pulse-pause mode [28]. Hence, an “apparent VHCF strength” rather than “intrinsic VHCF strength” [18] was employed in this work. Crack initiation regions (e.g., fish eye and GBF areas) of fracture specimens were observed directly by SEM and 3-Dimensional Confocal Microscope Phase Shift MicroXAM-3D (Vertical scan resolution: 0.01 nm, lateral resolution: 0.11 µm). Here, MicroXAM-3D can measure the surface profile of the fatigue fracture surface. In order to investigate the microstructural evolution in fish eye and GBF areas, the fracture surface was etched directly in a 2 vol% Nital solution (named “etched” fracture surface), subsequent observed by SEM. Furthermore, the focused ion beam (FIB, Helios Nanolab 600i) equipment was carried out to investigate the relationship between microstructure and NIICI.
results are listed in Table 1. It is found that the melting processes have little influence on the conventional mechanical properties of the B/M steels. The mechanical properties, especially impact toughness, are closely related to the BQ-P-T process parameters, which decide the final microstructure of the B/M steels [27]. The final microstructures of the steels after different BQ-P-T treatment have been characterized by SEM, TEM and XRD in our previous work [20]. The smelting process has little influence on the microstructures (see in Supplementary material). Based on our previous work [20], the significant difference of microstructure between the two BQP-T treatments lays in the size, volume fraction and morphology of retained austenite (RA): nanometer-sized filmy RA in Q200 specimens but micrometer-sized blocky RA in Q320 counterparts. Here, the B/M blocks are also characterized by EBSD in this present work, as shown in Fig. 2. The EBSD results companied with pole figure analyses show that the B/M microstructures in both Q200 and Q320 specimens maintain well the Kurdjumov-Sachs (K-S) orientation relationship with their parent austenite [27]. However, almost all of 24 variants exist within a single prior austenite grain (PAG) in Q200 specimens, whereas there are only 4 or 6 variants within an individual PAG in Q320 specimens. It means that more numbers and consequent smaller size of block are formed in Q200 specimens [27]. Through the interception line method, the measured average length of block in Q200 specimens is 9.1 ± 3.8 µm, which is much less than that (15.5 ± 3.6 µm) in Q320 counterparts. In summary, four specimens with various microstructure and inclusion sizes were designed in the present work: VQ200 specimen with large inclusion size but fine microstructure (refined B/M block and nano-sized filmy RA); EQ. 200 specimen with small inclusion size and fine microstructure; VQ320 specimen with large inclusion size and coarse microstructure; EQ. 320 specimen with small inclusion size but coarse microstructure. The mechanical, microstructural and inclusion properties of these four specimens are listed in Table 1.
3. Results 3.1. S-N characteristics Fig. 3a-d shows the S-N data obtained from VHCF testing of the four B/M steel specimens. The apparent mean fatigue strength [18] was determined by the conventional staircase method (i.e., up and down method, seen in supplementary material) by setting the run-out number of cycle near 109 [29]. The abbreviations in the upper right corner mean the crack initiation modes. The apparent mean fatigue strength (σw9) of the VQ200, EQ. 200, VQ320 and EQ. 320 specimens are 592 ± 47 MPa, 768 ± 13 MPa, 570 ± 44 MPa and 603 ± 13 MPa, respectively. The ratios of apparent mean fatigue strength to tensile strength (σw9/Rm) of the four variants are 0.42 ± 0.03, 0.52 ± 0.01, 0.43 ± 0.03 and 0.45 ± 0.01, respectively (Fig. 3f). It can be found that the inclusion size and microstructure have a
2.2. Fatigue testing method The geometry and dimensions of the smooth hour-glass type specimens for VHCF testing are shown in Ref [20]. The specimen geometry is dependent on the Young's modulus and density of steel, which
Fig. 2. EBSD observation showing the size and orientation of bainite/martensite blocks in (a) EQ. 200 and (b) EQ. 320 samples; bold black lines: misorientation, θ ≥ 15°, white lines: prior austenite grain boundaries, the insets showing {0 0 1} pole figure. 406
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Fig. 3. S-N data of the four B/M multiphase steels treated by various melting and BQ-P-T processes, (a) VQ200, (b) EQ. 200, (c) VQ320, (d) EQ. 320, and (e) the change of fatigue strength as well as (f) the ratio ofσw9/Rm for the four specimens; Sur: surface induced crack initiation, Sur (Inc): surface inclusion induced crack initiation, Int (Inc): interior inclusion induced crack initiation, Int (Non-inc): interior non-inclusion induced crack initiation (crack initiated from microstructure) [20].
3.2. Fracture surface observation
mutual and relative effect on the VHCF property of B/M steel, in other word, the effect of inclusion size on the fatigue strength strongly depends on the microstructure, and vice versa. For instance, the reduction of inclusion size can bring about 176 MPa increment of σw9 for the specimens with fine microstructure (VQ200 and EQ. 200), whereas the increment is only 33 MPa for the specimens with coarse microstructure (VQ320 and EQ. 320), as indicated by red arrows in Fig. 3e. Similarly, the refinement of microstructure has little influence on fatigue strength when the inclusion size is large, but the advantage of microstructural refinement becomes exciting when the inclusion size is reduced, as indicated by black arrows in Fig. 3e.
Based on fracture surface observation, the fatigue crack initiations are mostly induced by inclusions in the specimens with large inclusion size, e.g., 100% (11/11) in VQ200 specimen and 86.7% (13/15) in VQ320 variant. In contrast, no IICI is found in the specimens with small inclusion size due to electroslag remelting. But the fatigue cracks could initiate directly from interior microstructure when the microstructure is coarse, e.g., 13.3% (2/15) in VQ320 specimen and 92.3% (12/13) in EQ. 320 variant. Fig. 4a and b show the crack initiations are induced by surface inclusions or interior inclusions in VQ200 specimen. The GBF can be observed around the interior inclusions (Fig. 4b), but it is not evident 407
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Fig. 4. SEM images showing fractography of the VQ200 specimen, (a) surface inclusion (σa = 650 MPa, Nf = 8.75 × 105 cycles), (b) interior inclusion with GBF (σa = 575 MPa, Nf = 8.99 × 107 cycles).
Fig. 5. SEM images showing a “Non-inclusion induced crack initiation” fracture surface of the VQ320 specimen (σa = 600 MPa, Nf = 9.44 × 106 cycles) and (b) is the detailed observation of GBF; the cross in (b) showing the EDS gathering area, the inset in (b) showing the EDS spectrum.
Fig. 6. SEM images showing a “Non-inclusion induced crack initiation” fracture surface of the EQ. 320 specimen (σa = 650 MPa, Nf = 9.51 × 106 cycles) and (b) is the detailed observation of GBF.
which will be discussed in the following. Besides, only surface induced fracture is observed and no fatigue failure occurs at number of cycles exceeding 107 for EQ. 200 specimen [20], as shown in Fig. 3b. U. Krupp et al. [30] also reported the similar phenomenon and proposed that the absence of the failure at number of cycles exceeding 107 indicated that the early micro-cracks are nonpropagating cracks and permanently blocked by microstructural barriers. It is implied that, different from the other three specimens, neither interior inclusions nor microstructure could induce the crack initiation during VHCF regime for EQ. 200 specimen.
around the surface inclusions (Fig. 4a). Fig. 5 and Fig. 6 show the fracture surfaces of NIICI in the VQ320 and EQ. 320 samples, respectively. The fish-eye with GBF is also formed in the non-inclusion induced fracture surface. However, the detailed EDS analysis indicates that no inclusions exist around the GBF and the composition in the noninclusion induced GBF is similar to that in the matrix, which confirms that the crack initiation is induced by interior microstructure rather than inclusions [20]. The SEM observation in Fig. 6b shows the presence of numerous submicron-sized facets and the trace of severe plastic deformation within GBF. The severe plastic deformation might come from the rubbing of two fracture surface or the strain localization, or both. The former is generated after the formation of cracks, whereas the later occurs before crack initiation and may be the cause of NIICI, 408
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3.3. Stress intensity factor range The parameters of characteristic area can be obtained from fracture surface observation. Then, the value of stress intensity factor (SIF) range at the periphery of inclusion (ΔKInc), GBF around inclusion (ΔKIG) and GBF without inclusion (ΔKNG) are calculated based on the area model [31]:
∆K = 0.5σa π area
(1)
where σa is stress amplitude, area is the projected area of inclusion and GBF on the plane perpendicular to tensile stress and area is used to characterized the size of inclusion and GBF [32]. For surface inclusion or defect, the coefficient of 0.5 is replaced by 0.65 in the calculation [32]. Fig. 7 shows the calculated value of ΔKInc, ΔKIG and ΔKNG versus fatigue life. The values of ΔKInc are between 2.0 and 4.5 MPa√m, with the value of areaInc ranging from 10 µm to 83 µm (Fig. 1a). The values of ΔKIG are between 3.3 and 6.6 MPa√m, while those of ΔKNG are relatively small, varying from 2.9 to 5.0 MPa√m. The important point was that the values of ΔKIG and ΔKNG almost keep constant at average of 4.8 MPa√m and 3.8 MPa√m under the failure cycles between 106 and 109, whereas the values of ΔKInc display a slightly decreasing tendency with respect to failure life. This confirms the correlation between inclusion size and fatigue life.
Fig. 7. Stress intensity factor range of inclusion (Inc), GBF with inclusion (IGBF) and GBF without inclusion (N-GBF) of the VQ200, VQ320 and EQ. 320 specimens.
Fig. 8. MicroXAM-3D observation of an “inclusions induced crack initiation” fracture surface in VQ320 specimen (σa = 600 MPa, Nf = 1.47 × 108 cycles), (a) the optical image, the inset in (a) showing the EDS spectrum, (b) 3D morphology image, (c) the sectional profile of the labeled line in (b), the bottom is an enlarged window showing the detailed observation of GBF. 409
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Fig. 9. MicroXAM-3D observation of an “Non-inclusions induced crack initiation” fracture surface in EQ. 320 specimen (σa = 650 MPa, Nf = 9.51 × 106 cycles), (a) the optical image, (b) 3D morphology image, (c) the sectional profile of the labeled line in (b), the bottom is an enlarged window showing the detailed observation of GBF. (Corresponding SEM observation is shown in Fig. 6).
strength strongly depends on the microstructure, and vice versa. For the refined microstructure, the reduction of inclusion size could bring an exciting increment of fatigue strength, but the efficacy is weak for the coarse microstructure. It is because, although the IICI is eliminated due to the reduction of inclusion size, the NIICI would become the dominant fracture mode when the microstructure is coarse (e.g., in EQ. 320 specimen). In other word, there is a competition process between nonmetallic inclusion and microstructure to determining the VHCF properties of B/M steels [34]. Hence, it is essential to understand the mechanisms of IICI and NIICI to explore the crucial factors in controlling the fracture modes. Since the mechanism of IICI has been reported in many literatures [22–24], we will focus on the mechanism of NIICI in the following. Through the surface profiles analysis by MicroXAM-3D, we find that, at macroscopic scale, the surface of inclusion induced GBF (IG) is normal to the loading direction (Fig. 8), but the non-inclusion induced GBF (NG) is along close to the plane of maximum shear stress (Figs. 9 and 10). The surface profiles of NG have a strong similarity with the stage I crack propagation, which is identified in single crystals with crack growing in a plane oriented at an angle about 45° against the loading direction [35]. Stage I crack propagation is also typical of the early propagation of micro-cracks in surface grains as shown by Forsyth [36]. This stage is sensitive to microstructure and related to slip/strain localization in individual grains [35]. Owing to the irreversible dislocation slip, the persistent slip bands (PSBs) accompanied with
3.4. Fracture mechanisms There are two dominant interior crack initiation modes for B/M steels within VHCF regimes, namely, IICI and NIICI modes. In order to investigate the difference between the fracture mechanisms of two crack initiation modes, the MicroXAM-3D was employed to measure the surface profile of the fatigue fracture surface. Fig. 8 shows the IICI fracture surface profile in the VQ320 specimen. The failure is caused by interior inclusion, and the GBF with fish eye can be observed around the inclusion (the GBF is also called optically dark area, ODA, Fig. 8a). As shown by the profile (Fig. 8c), the surfaces of GBF and fisheye are all perpendicular to the loading direction, which indicates that the early cracks in GBF are formed directly due to inclusions and then grow along the plane of maximum tensile stress [33]. Figs. 9 and 10 show the NIICI fracture surface profiles of the EQ. 320 specimen. The surface profiles show the GBF actually contains a micro-facet with size of tens of micrometer. This facet has an angle of about 40° against the plane of maximum tensile stress, namely, close to the plane of maximum shear stress. It could be inferred that, unlike inclusion induced cracks, the early propagation of the non-inclusion induced cracks within GBF is controlled by maximum shear stress [33]. 4. Discussions As indicated in Fig. 3, the effect of inclusion size on the fatigue 410
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Fig. 10. MicroXAM-3D observation of an “Non-inclusions induced crack initiation” fracture surface in EQ. 320 specimen (σa = 675 MPa, Nf = 1.89 × 106 cycles), (a) the optical image, (b) 3D morphology image, (c) the sectional profile of the labeled line in (b), the bottom is an enlarged window showing the detailed observation of GBF.
micro-facet through continued cyclic loading. Another question arises: since the orientation of bainite/martensite lath is on the whole random (in whether refined or coarse microstructure), the conditions of strain localization could be fulfilled for any lath along 45° oriented {110} plane against the loading direction [38]. However, we observe the NIICI only in the Q320 specimens with coarse microstructure. In order to explain this phenomenon, we observe the NIICI fracture surface on a relatively large scale. Fig. 12 shows a typical NIICI fracture surfaces before and after etching of an EQ. 320 specimen failed at Nf = 7.62 × 107 with the stress amplitude σa = 600 MPa. It can be found that the region of fish eye envelops a GBF (equivalent length 46.2 µm), but no inclusion is found around GBF (Fig. 12a and b). After etching, the micro-facet is obviously observed, whose size is smaller than that of GBF (Fig. 12c). Through backscattered electron image (Fig. 12d), we find that the micro-facet basically runs through the prior austenite grain (PAG) and directly impinges on the PAG boundaries. And the micro-voids and micro-cracks are observed along the meeting place of micro-facet and PAG boundaries. It implies that the orientation and dimension of microfacet and the state of PAG boundaries are crucial factors in determining whether the strain localization could develop the NIICI crack. Based on the discussion above, the mechanism of NIICI in VHCF regime for the coarse B/M microstructure is suggested:
extrusion and intrusions lead to the surface relief which acts as stress raiser due to its notch effect in the continuation of the cyclic deformation and as a consequence can lead to crack initiation [37]. In the present work, it is interesting that the Stage I crack propagation is found in the interior microstructure, but in which case, the relief at free surface cannot form. Then, a question arises that how the crack initiates from interior microstructure. We carried out the focused ion beam (FIB) sample cut from the micro-facet to observe the correlation of microstructural characteristic and crack initiation origin, as shown in Fig. 11. It can be found that the profile of FIB sample (Fig. 11c) is consistent with the results of MicroXAM-3D (Figs. 7c and 8c), which confirms that the micro-facet is along the plane of maximum shear stress. A sheaf of parallel bainite/ martensite lath is observed beneath the micro-facet, and the micro-facet lies parallel to the lath longitudinal direction. Furthermore, the selected area diffraction (SAD) patterns show that the {110} planes of the B/M lath is also along the direction of maximum shear stress (Fig. 11d). L. Morsdorf et al. [38] studied the plasticity of lath martensite by EBSD, digital image correlation (DIC) and electron channeling contrast imaging (ECCI) and found that the slip in the lath preferred to activate along the 45° oriented {110} plane against the loading direction, where the shear stress is highest. Hence, the localized slip/strain could firstly occur along ~ 45° oriented laths, which could result in the formation of 411
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Fig. 11. Focused ion beam (FIB) observation showing the correlation of microstructural characteristic and crack initiation of VQ320 sample (σa = 600 MPa, Nf = 4.38 × 107), (a) SEM image showing the GBF area, (b) SEM image showing the location of FIB cut, Pt coating was to protect the fracture surface, (c) SEM image showing the profile of fracture surface, (d) TEM image and SAD showing the microstructure beneath GBF and the orientation of B/M lath.
Fig. 12. SEM image showing typical “non-inclusion induced crack initiation” fracture surfaces before (a, b) and after etching (c, d) of an EQ. 320 specimen (σa = 600 MPa, Nf = 7.62 × 107), PAGB: prior austenite grain boundary, RA: retained austenite. 412
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Fig. 13. The schematic representation of NIICI mechanism in coarse B/M microstructure, the occurrence of NIICI progresses according to the order of (a) → (b) → (c) → (d), PAGB: prior austenite grain boundary, RA: retained austenite, M: newly formed martensite from RA under cyclic loading.
Furthermore, the nano-sized filmy RA located between B/M lathes has high stability and also acts as strong barrier to slip movement and crack propagation due to the dislocation absorption and transformation induced plasticity effects [45,46]. In this case, the early localized slip bands are permanently blocked by these microstructural barriers, and consequently the NIICI cannot occur in the Q200 specimen with refined microstructure.
The first stage (Fig. 13a): under cyclic loading, slip localization prefers to occur along some certain oriented crystallographic planes [39], e.g., the {110} planes oriented in 45° against the loading direction (as shown in Fig. 11d). The bainite and martensite in low carbon steels have a three-level hierarchy in its morphology, i.e. packet, block and lath [40]. The packet is a group of laths with a common habit plane with respect to austenite. Each packet is divided into blocks, which contain the parallel laths with the same crystallographic orientation and is also called as a co-variant packet. Since the micro-facet is parallel to the longitudinal direction of the B/M laths with the same {110} plane oriented in 45° against the loading direction (Fig. 11c), it is deduced that the length of blocks governs the slip localization at the early stage of fatigue for B/M steels [41]. This was also confirmed by Mueller et al. [42] through building the constitutive relationship between block size and fatigue limit in nano-structured bainitic steels. The second stage (Fig. 13b): with increasing cycles, in case that the B/M block is coarse and grows through the PAG (e.g., the B/M block in Q320 specimen, Fig. 2b), the localized slip bands could directly impinge on the packet or PAG boundaries without any microstructural barrier. At the third stage (Fig. 13c), whether or not the localized slip bands could grow across the PAG boundaries into the adjacent PAG depends on the state of PAG boundaries and the critical cyclic stress for dislocation moving [43]. Since most of blocky RA is located nearby the PAG boundaries in EQ. 320 specimen [20], the strength of PAG would be reduced. Furthermore, the blocky RA along PAG boundaries transforms easily to brittle martensite due to localized plastic deformation under the impinging of slip bands on PAG boundaries [20]. The brittle martensite is prone to induce the formation of cleavage micro-cracks along the PAG boundaries (as shown in Fig. 12d). At the final stage (Fig. 13d), the micro-cracks and voids along the PAG boundaries act as stress raisers and help the localized slip bands grow across the PAG boundaries, and then the occurrence of NIICI becomes possible. The schematic representation of NIICI mechanism in coarse B/M microstructure is shown in Fig. 13. In contrast, the orientation distribution of refined B/M blocks are intersecting in Q200 sample (as shown in Fig. 2a), which supplies more boundary barriers to restricting the localized slip progression [44].
5. Conclusions In this work, four B/M steel specimens with various microstructure and inclusion size were prepared by controlling the melting and BQ-P-T processes, i.e., VQ200 specimen with large inclusion size (maximum size: ~ 98 µm) but fine microstructure (refined B/M block and nanosized filmy RA); EQ. 200 specimen with small inclusion size (maximum size: ~ 15 µm) and fine microstructure; VQ320 specimen with large inclusion size and coarse microstructure; EQ. 320 specimen with small inclusion size but coarse microstructure. The effect of inclusions and microstructure on the very high cycle fatigue behaviors of B/M multiphase steel was investigated through S-N curves testing, fracture surface observation, failure mechanism and crack initiation analyses. In particular, the mechanism of “non-inclusion induced crack initiations” was investigated. The main conclusions are described as follows: (1) The inclusion size and microstructure have a mutual and relative effect on the VHCF property of B/M steel, i.e., the effect of inclusion size on the fatigue strength strongly depends on the microstructure, and vice versa. (2) The “inclusion induced crack initiation” (IICI) is a leading failure mode within VHCF regimes for the specimens with large inclusion size, while we observe the “non-inclusion induced crack initiations” (NIICI) only in the specimens whose microstructure is coarse. (3) At macroscopic scale, the surface of inclusion induced GBF is normal to the loading direction, but the non-inclusion induced GBF is along close to the plane of maximum shear stress. (4) The NIICI micro-facet is parallel to the longitudinal direction of the B/M laths with {110} plane oriented in 45° against the loading 413
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direction，which is related to the strain localization during cyclic loading. And the micro-voids and micro-cracks are observed along the meeting place of NIICI micro-facet and PAG boundaries. (5) The orientation and size of coarse B/M block and blocky RA are crucial microstructural factors in facilitating the formation of NIICI.
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Acknowledgment This research is supported by National Natural Science Foundation of China (No. 51771014 and U1760104), the Fundamental Research Funds for the Central Universities (No.2017JBM053), and Natural Science Foundation of Beijing (2172047). The authors would also like to express their appreciation to Central Iron & Steel Research Institute for supporting this study and providing ultrasonic fatigue testing equipment (SHIMADZU USF-2000). Declarations of interest None. Appendix A. Supporting information Supplementary data associated with this article can be found in the online version at doi:10.1016/j.msea.2018.10.073. References [1] T. Sakai, Y. Sato, N. Oguma, Characteristic S–N properties of high‐carbon–chromium‐bearing steel under axial loading in long‐life fatigue, Fatigue Fract. Eng. Mater. 25 (8‐9) (2002) 765–773. [2] K. Shiozawa, L. Lu, Very high‐cycle fatigue behaviour of shot‐peened high‐carbon–chromium bearing steel, Fatigue Fract. Eng. Mater. 25 (8‐9) (2002) 813–822. [3] C. Bathias, C.P. Paris, Gigacycle Fatigue in Mechanical Practice, Taylor & Francis, 2004. [4] B. Pyttel, D. Schwerdt, C. Berger, Very high cycle fatigue–is there a fatigue limit? Int. J. Fatigue 33 (1) (2011) 49–58. [5] Q. Jiang, C. Sun, X. Liu, Y. Hong, Very-high-cycle fatigue behavior of a structural steel with and without induced surface defects, Int. J. Fatigue 93 (2016) 352–362. [6] T. Sakai, A. Nakagawa, N. Oguma, Y. Nakamura, A. Ueno, S. Kikuchi, A. Sakaida, A review on fatigue fracture modes of structural metallic materials in very high cycle regime, Int. J. Fatigue 93 (2016) 339–351. [7] K. Shiozawa, Y. Morii, S. Nishino, L. Lu, Subsurface crack initiation and propagation mechanism in high-strength steel in a very high cycle fatigue regime, Int. J. Fatigue 28 (2006) 1521–1532. [8] S.X. Li, Effects of inclusions on very high cycle fatigue properties of high strength steels, Int Mater. Rev. 57 (2) (2012) 92–114. [9] A. Pineau, S. Forest, Effects of inclusions on the very high cycle fatigue behaviour of steels, Fatigue Fract. Eng. Mater. 40 (11) (2017) 1694–1707. [10] U. Karr, R. Schuller, M. Fitzka, B. Schönbauer, D. Tran, B. Pennings, H. Mayer, Influence of inclusion type on the very high cycle fatigue properties of 18Ni maraging steel, J. Mater. Sci. 52 (10) (2017) 5954–5967. [11] D.L. McDowell, F.P.E. Dunne, Microstructure-sensitive computational modeling of fatigue crack formation, Int J. Fatigue 32 (9) (2010) 1521–1542. [12] G.M. Castelluccio, W.D. Musinski, D.L. McDowell, Recent developments in assessing microstructure-sensitive early stage fatigue of polycrystals, Curr. Opin. Solid State Mater. 18 (4) (2014) 180–187. [13] P. Zhao, G. Gao, R.D.K. Misra, B. Bai, Effect of microstructure on the very high cycle fatigue behavior of a bainite/martensite multiphase steel, Mater. Sci. Eng. A Struct. 630 (2015) 1–7. [14] W. Li, T. Sakai, M. Wakita, S. Mimura, Influence of microstructure and surface defect on very high cycle fatigue properties of clean spring steel, Int. J. Fatigue 60 (2014) 48–56. [15] G. Chai, The formation of subsurface non-defect fatigue crack origins, Int. J. Fatigue 28 (11) (2006) 1533–1539. [16] G. Chai, N. Zhou, Study of crack initiation or damage in very high cycle fatigue using ultrasonic fatigue test and microstructure analysis, Ultrasonics 53 (8) (2013) 1406–1411. [17] J. Zhang, S. Li, Z. Yang, Z. Yang, G. Li, W. Hui, Y. Weng, Influence of inclusion size on fatigue behavior of high strength steels in the gigacycle fatigue regime, Int. J. Fatigue 29 (4) (2007) 765–771. [18] Y. Yu, J. Gu, L. Xu, F. Shou, B. Bai, Y. Liu, Very high cycle fatigue behaviors of Mn–Si–Cr series bainite/martensite dual phase steels, Mater. Des. 31 (6) (2010)
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