Microstructural Influences on Very-High-Cycle Fatigue-Crack Initiation in Ti-6246
- First Online:
- Cite this article as:
- Szczepanski, C., Jha, S., Larsen, J. et al. Metall and Mat Trans A (2008) 39: 2841. doi:10.1007/s11661-008-9633-z
- 1.2k Downloads
The fatigue behavior of an alpha + beta titanium alloy, Ti-6Al-2Sn-4Zr-6Mo, has been characterized in the very-high-cycle fatigue (VHCF) regime using ultrasonic-fatigue (20 kHz) techniques. Stress levels (σmax) of 40 to 60 pct of the yield strength of this alloy have been examined. Fatigue lifetimes in the range of 106 to 109 cycles are observed, and fatigue cracks initiate from both surface and subsurface sites. This study examines the mechanisms of fatigue-crack formation by quantifying critical microstructural features observed in the fatigue-crack initiation region. The fracture surface near the fatigue-crack-initiation site was crystallographic in nature. Facets, which result from the fracture of primary alpha (αp) grains, are associated with the crack-initiation process. The αp grains that form facets are typically larger in size than average. The spatial distribution of αp grains relative to each other observed near the initiation site did not correlate with fatigue life. Furthermore, the spatial distribution of αp grains did not provide a suitable means for discerning crack-initiation sites from randomly selected nominal areas. Stereofractography measurements have shown that the facets observed at or near the initiation sites are oriented for high shear stress; i.e., they are oriented close to 45 deg with respect to the loading axis. Furthermore, a large majority of the grains and laths near the site of crack initiation are preferentially oriented for either basal or prism slip, suggesting that regions where αp grains and α laths have similar crystallographic orientations favor crack initiation. Microtextured regions with favorable and similar orientations of αp grains and the lath α are believed to promote cyclic-strain accumulation by basal and prism slip. Orientation imaging microscopy (OIM) indicates that these facets form on the basal plane of αp grains. The absence of a significant role of spatial clustering of αp grains, coupled with the observation of regions of microtexture on the order of 300 to 500 μm supports the idea that variability in fatigue life in the very-high-cycle fatigue regime results from the variability in the nature (intensity, coherence, and size) of these microtextured regions.
The study of very-high-cycle fatigue (VHCF) behavior is attracting increased interest from industry because many components in structural applications, such as automobile cylinder heads, engine blocks, and turbine engines, will accumulate 108 to 1010 cycles in service. The conventional approach of designing components to a fatigue limit is not applicable in VHCF because fatigue failures have been observed below the conventional fatigue limit. This has led some researchers to propose a modified stress-life (SN) curve where surface-initiated fatigue failures are observed at high stresses, and subsurface fatigue-crack initiation is observed at very long lifetimes below the conventional fatigue limit. Mughrabi explained failures below the conventional fatigue limit in terms of fundamental physical mechanisms of fatigue-damage accumulation by suggesting that, even though macroscopic strain in VHCF is below the persistent slip band threshold, slip irreversibility can still accumulate and lead to failure. This agrees with the work of Lukáš and Kunz who argue that in the VHCF regime the applied strain is nominally elastic, and only localized plastic deformation will accumulate at specific microstructural locations. Thus, as compared to low cycle fatigue where the majority of grains accommodate some plastic deformation, in the longer lifetime regimes of high-cycle fatigue (HCF) and VHCF (i.e., ≥106 cycles), it is likely that specific microstructural configurations can be associated with more rapid local fatigue-damage accumulation. Such fatigue-critical microstructural neighborhoods have been defined based on grain size, spatial orientation,[9,10] proximity to the specimen surface,[3,6] and crystallographic orientation.[7,17] The current work examines which of these factors, if any, contribute to the process of fatigue-crack initiation in the α + β titanium alloy Ti-6Al-2Sn-4Zr-6Mo.
Much of the literature on fatigue-crack initiation in α + β titanium alloys focuses on analysis of the fracture surfaces to determine the deformation and possible strain-accumulation mechanisms that are responsible for crack initiation. In materials that do not contain inclusions or porosity, fatigue cracks tend to initiate in locations where the local microstructure promotes the accumulation of irreversible slip. Hall’s review of fatigue-crack initiation in α + β titanium alloys established that fatigue damage typically accumulates in the alpha phase. Further, the length scale of the deformation will vary depending on the microstructure and processing conditions. Fatigue-critical microstructural features have been identified as individual αp grains, α colonies, prior β grains, or regions of similarly oriented α grains. Mahajan and Margolin found that, in Ti-6246, fatigue cracks initiated in large alpha grains or in areas where a number of alpha grains was clustered, which presumably increased the slip length. They hypothesized that a likely method for improving the fatigue resistance of this alloy is to increase the spacing between alpha grains or to refine the αp grain size to limit slip transmission between αp grains or plastic deformation within grains, respectively. Researchers have investigated clustering both in a statistical and a micromechanistic approach. Chandran and Jha determined that αp clustering is the fatigue-critical microstructural feature in Ti-10V-2Fe-3Al in one population of failures, and they were able to model this using Poisson defect statistics. In Waspaloy (Inco Alloys, Int., Huntington, WV), Davidson et al. found that the crack-initiation sites are associated with clusters of similarly-oriented grains, which they termed “supergrains.” They suggested that supergrains were more susceptible to fatigue-crack initiation because localized deformation in one grain could be more easily accommodated in adjacent grains due to their similar crystallographic orientation. Regions of material with similar crystallographic orientation are commonly observed to initiate fatigue cracks in titanium alloys, as well.[7,8,12,17] Bieler and Semiatin established that the presence of these microtextured regions results from local heterogeneities in deformation during thermomechanical processing.
Macroscopic textures, resulting from the processing history of titanium alloys, are known to influence mean fatigue properties.[14,15] However, in the regime of VHCF, damage accumulation will only occur if the local microstructure is suitable for irreversible slip to accumulate. In other words, microstructural heterogeneity is thought to cause scatter in fatigue lifetimes. Evidence of such behavior has been observed by a number of researchers who have investigated the effect of local texture on fatigue-crack initiation[7,17] and propagation in titanium alloys. In Ti-6242, Sinha et al. have found that dwell-fatigue loading, in which specimens are held at a static load as part of every fatigue cycle, leads to crack formation in microtextured regions suitable for basal slip that are surrounded by regions of material oriented for prism slip. Multiple cracks initiated, and while the dominant crack was not necessarily the first to initiate, they did grow out of the largest microtextured region. In the work of LeBiavant et al., numerous cracks were found to initiate in macrozones where the majority of αp grains were oriented for basal or prism slip. Each of these studies aimed to understand the contribution of heterogeneous texture to the initiation and growth of fatigue cracks in the HCF regime. The objective of the current work is to establish which microstructural features and configurations cause fatigue-crack initiation in the VHCF regime to better understand the mechanisms of damage accumulation and fatigue-crack initiation.
2 Materials and experimental procedures
Specimens were cut circumferentially from a forged pancake that was processed to simulate forging conditions in an actual component. Cylindrical specimen blanks were cut from this material, and grip ends of Ti-6Al-4V rod were inertia welded onto the specimen blanks. Final machining of cylindrical specimens was completed by low-stress grinding to minimize compressive residual stresses. All specimens were then electropolished to remove the remaining surface compressive residual stresses. The gage was 4 mm in diameter and 12 mm in length. Axial fatigue testing was completed using ultrasonic-fatigue techniques detailed elsewhere.[20–22] Specimens were designed such that their resonant frequency was approximately 20 kHz.
Observation of the fracture surfaces was completed using two scanning electron microscopes (SEMs): a Leica Cambridge S360FE microscope (Leica Microsystems GmbH, Wetzlar, Germany) operating at a probe current of 100 pA and accelerating voltage of 20 kV, and a PHILIPS1 XL30FEG operating at 10 kV with a probe current of approximately 2 nA. Orientation imaging microscopy (OIM) was performed on polished sections of the fractured fatigue specimens using these two microscopes with detectors manufactured by EDAX-TSL (Ametek, Mahwah, NJ). For the OIM investigations, the probe current and accelerating voltage were approximately 10 nA and 20 kV, respectively.
To gain insight into the process of crack initiation, the spatial orientations of facets were measured with respect to the loading axis using stereography. Images used to produce three-dimensional (3-D) reconstructions of fracture surfaces were acquired in the SEM by centering the feature of interest in the field of view at the eucentric height. Images were then acquired at tilt angles of 0 and 6 deg with respect to the stage normal position for each facet. The technique and procedure for generating stereo pairs was validated using MeX commercial software (Alicona Imaging GmbH, Grambach, Graz, Austria) on known geometries, such as Vickers microhardness indents.
3 Results and discussion
3.1 Fatigue Lifetime
Additionally, there was an increased likelihood of subsurface-crack initiation as the stress level is decreased, which was also consistent across the testing frequencies. At ultrasonic frequencies, surface-crack initiation was observed approximately 80 pct of the time for the highest stress level investigated (700 MPa), while at the lowest stress level (550 MPa), only 20 pct of the specimens failed from a surface-initiated fatigue crack. Specimens that failed by surface-crack initiation typically tended to have shorter lifetimes than those that failed from subsurface-crack initiation. However, in a few cases, surface-initiated failures were observed to have similar lifetimes as specimens that failed from subsurface-crack initiation sites, consistent with other investigations of fatigue in α + β titanium alloys.
Subsurface-crack initiation has been observed in a number of fatigue studies on α + β titanium alloys.[10,23,32–35] The competition between surface and subsurface crack-initiation sites has alternately been attributed to the specimen surface to volume ratio, the presence of compressive residual stresses on the surface, environmental effects, and the relative ease with which grains can deform at a free surface as compared to the specimen interior.
In the regime of VHCF, only certain microstructural neighborhoods are susceptible to fatigue-damage accumulation that leads to crack initiation, and these microstructural configurations do not necessarily exist at the specimen surface. This idea was first put forth by Mughrabi, and Jha and Larsen subsequently proposed that different microstructural neighborhoods may be in competition with one another to initiate the dominant fatigue crack. In the current work, it appears that there may be a number of different microstructural regions in each sample that may initiate a fatigue crack; however, only the site that accumulates fatigue damage most rapidly will ultimately lead to fatigue-crack initiation. The fact that three characteristic crack-initiation sites have been observed at the same stress level indicates that there is competition between different microstructural neighborhoods for crack initiation. The evidence suggests that if there is not a suitable group of αp grains near the specimen surface, then a crack will initiate from the specimen interior, and the presence of different microstructural neighborhoods in the interior will cause a subsurface-initiation site with either isolated facets or with a macroscopically planar region. The occurrence of αp facets at the crack-initiation sites indicates that αp grains are influential in accumulating fatigue damage and subsequent fatigue-crack initiation. Similar observations led to speculation that specific spatial and crystallographic orientations of αp grains are required for fatigue cracks to initiate.
3.3 Spatial Clustering of αp Grains
A number of observations can be drawn from these plots. First, the area fraction of αp facets typically decreases as material farther from the crack-initiation site is considered. As the facet-formation process is suspected to result from fracture of αp grains on crystallographic planes favorable for slip, the apparent lack of faceting for large crack sizes is expected because long cracks are not expected to propagate in a crystallographic manner. The local microstructure is known to influence crack initiation and small-crack growth;[11,17,29] therefore, more facets are expected to form in these regions than in regions where long crack-growth conditions apply.
The results suggest that there is no critical value of spatial clustering of αp that must be met for cracks to initiate. In some cases, the αp facets are much more dispersed at the site of crack initiation than the αp grains are in the general microstructure, while in other samples, the degree of faceting at the initiation site is markedly higher than the average αp-grain distribution. Thus, the spatial distribution of αp grains alone does not determine where cracks will initiate in the general microstructure. Additionally, the lifetimes of specimens do not correlate with the spatial distribution of αp at the crack-initiation sites. Assuming that the majority of αp grains that the fatigue-crack encounters fail by faceted fracture, this suggests that the area fraction of faceting is not indicative of the rate of fatigue-damage accumulation or the degree of difficulty to initiate a fatigue crack in these microstructural regions. As the plots in Figures 8(a) and (b) illustrate, there is no discernable difference in the degree of faceting between surface- and subsurface-crack initiations. It is apparent that, in general, more faceting is observed at the crack-initiation sites for higher σmax, as shown in Figure 8(c) and (d). If faceting resulted from crystallographic propagation of short fatigue cracks, this trend would be expected to be reversed. Therefore, this finding suggests that the higher resolved stresses on the basal and prism slip systems in unfavorably oriented grains may cause slip to be activated in grains that would not otherwise contribute to the damage-accumulation process. As a result, a larger volume of material could contribute to the crack-initiation process at higher stress levels.
3.4 Spatial Orientation of Facets
Spatial Orientation of Facets with Respect to Fracture Surface (in Deg)
3.5 Crystallographic Orientation of Crack Initiating Region
Orientation of Facets (in Deg)
Spatial orientation (with respect to tensile axis)
Crystallographic orientation ( with respect to tensile axis)
Thus, it appears that the presence of a preferred texture suitable for basal and prism slip is a necessary, but not a sufficient, condition for crack initiation. In these regions of preferred texture, it is clear that the majority of the α-phase material shares the same crystallographic orientation. However, the strength of the preferred texture will affect the resistance to fatigue-crack initiation. In other words, a weakly textured region of material will not promote crack initiation and propagation as successfully as a suitably oriented strongly textured region might. It has been established that regions of material with preferred texture for basal and prism slip are found at the crack-initiation site. However, quantification of these variations in texture intensity has not been completed, and this will be examined in future work.
Thus far, the texture of the α phase, i.e., the lath α and the αp grains, has been considered as one. However, the texture of the lath α and equiaxed αp phases should be considered separately because the correlation in the textures of these phases will affect the ease of slip transmission between grains. The lath α maintain a Burgers orientation relationship with the β phase, whereby the basal (0001) plane of the α phase is parallel to the (110) plane of the β phase. Due to the spherodization process which forms the αp grains, it is generally assumed that equiaxed αp are not crystallographically related to the prior β-grain orientation. However, Woodfield et al. note that strains on the order of ε = 1 are required to dynamically spheroidize the α laths to form αp grains, but this strain still may not be high enough to cause recrystallization of the α lath material as they spheroidize. Therefore, the αp grains may, in fact, share a crystallographic-orientation relationship with the prior β, and hence, the lath α and slip transmission across phases may not be hindered by the grain boundaries. Quantification of these individual texture components remains as future work.
A similar explanation for crack initiation has been postulated in the work of LeBiavant et al. on the microstructural influences of bending fatigue in Ti-6Al-4V. They observed that cracks initiate in macrozones, which are defined as regions of material in which the majority of αp grains have a similar orientation. In the macrozone that initiated the dominant fatigue crack, multiple microcracks were observed. They postulated that the dominant fatigue crack grew out of these macrozones due to microcrack coalescence. In the current work, there is no evidence of slip offsets or microcracks in αp grains, aside from those on the fracture surface. Fatigue damage is believed to accumulate throughout these similarly oriented regions of material, and slip activity in neighboring grains may facilitate crack initiation and growth. However, no evidence of other cracks is observed; thus, crack coalescence is not believed to significantly affect crack-growth rates. This difference may be due to the fact that LeBiavant et al. performed their tests at 800 MPa in a material with a yield strength of 850 MPa, while the current investigation focused on the fatigue behavior at much lower stresses in the range of 0.4 to 0.6 of σYS, where it is unlikely that slip activity would intensify to such an extent that fatigue cracks would initiate in other αp grains.
The proposed mechanism can be generalized to explain the process of fatigue-crack initiation for the three types of crack-initiation sites that were observed. In the case of surface-crack initiation, faceted fracture of αp grains is still observed, and the mechanism is essentially the same. However, the size of the faceted region is smaller for surface failures than it is for subsurface crack-initiation sites. Therefore, the same mechanism is believed to be applicable, but fewer αp grains are required to accumulate cyclic strain because surface-connected cracks can be smaller than subsurface cracks and still propagate at the same rate due to differences in stress intensity and environmental contributions at these sites. In subsurface macroscopically planar failures, the whole faceted region has the same spatial orientation with respect to the loading axis. This indicates that the αp grains and lath-α colonies within the crack-initiation region all have a similar crystallographic orientation. The noticeable feature of these initiation sites is that they appear to form from the faceted fracture of transformed beta regions of material. There are clearly αp grains in that area; however, the slip planes of αp grains and transformed beta regions of material must be aligned such that the grain boundaries between these phases are not distinguishable on the fracture surface. For these types of failures, it is still thought that slip must accumulate in a number of αp grains or transformed beta regions in order to initiate a fatigue crack.
Ultrasonic fatigue at a load ratio of 0.05 and stresses (σmax) in the range of 500 to 700 MPa has been shown to produce failures that display similar trends in lifetime and fractographic appearance with respect to conventional-frequency fatigue. Cracks initiate in larger than average-sized αp grains by strain localization resulting from basal and prism slip. The initiation process leads to microcrack formation in the αp grains parallel to the basal plane. Pseudocleavage is not operative because facets are oriented for slip on rational crystallographic planes. Cracks typically form in αp grains with similar orientations that favor slip on basal planes in which the basal planes are only slightly misaligned in neighboring grains. These microstructural configurations appear to be found more often in textured regions. It is believed that this allows for relatively easy slip transmission between neighboring αp grains. However, spatial clustering of αp grains has been proven not to be the distinguishing microstructural feature at crack-initiation sites. Crack-initiation sites have a preferred texture, and it is observed that a majority of the grains near the site are oriented for easy basal or prism 〈a〉 type slip. Analysis of the fatigue-lifetime data using probabilistic models of fatigue-crack initiation will be presented in a forthcoming article.
PHILIPS is a trademark of FEI, Hillsboro, OR.
The authors thank the AFOSR Metallic Materials Program (Project No. F49620-03-1-0069) for financial support. One of the authors (CJS) thanks the STEP program at the AFRL Materials and Manufacturing Directorate for funding. The authors also thank C. Torbet, University of Michigan, for technical assistance.