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A study of cyclic fatigue, damage initiation, damage propagation, and fracture of welded titanium alloy plate
Materials Science and Engineering A 527 (2010) 6649–6659
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Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
A study of cyclic fatigue, damage initiation, damage propagation, and fracture of welded titanium alloy plate T.S. Srivatsan a,∗ , U. Bathini b , A. Patnaik b , T. Quick c a b c
Department of Mechanical Engineering, The University of Akron, Akron, OH 44325-9303, USA Department of Civil Engineering, The University of Akron, Akron, OH 44325-9303, USA Department of Geology, The University of Akron, Akron, OH 44325, USA
a r t i c l e
i n f o
Article history: Received 19 May 2010 Accepted 28 June 2010
Keywords: Welded titanium alloy Damage initiation Damage propagation Fatigue life Microstructure Fracture
a b s t r a c t In this paper, the influence of test specimen orientation and microstructure on cyclic stress-amplitude controlled fatigue response, damage initiation, damage propagation and fracture behavior of samples taken from a welded plate of titanium alloy is presented and discussed. Test specimens from the chosen alloy were prepared from an as-welded plate of the material with the stress axis both parallel (longitudinal) and perpendicular (transverse) to the deformed (rolling) direction of the plate. The test specimens were cyclically deformed at different values of maximum stress at a constant load ratio of 0.1, and the resultant cycles-to-failure was recorded. The fracture surfaces of the deformed and failed test specimens were examined in a scanning electron microscope to establish the macroscopic fracture mode, the intrinsic features on the fatigue fracture surface and the role of applied stress-microstructural feature interactions in establishing the microscopic mechanisms governing failure. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The ever increasing scarcity of resources coupled with their growing expense demands both attention and action towards a reduction in the consumption of energy when dealing with the transportation of passengers and goods. It is here that the aerospace sector comes into to picture in playing an important and key role with specific respect to the development, emergence, application and commercialization of new and improved materials. In direct comparison with the land-based transportation systems, the much lower system quantities coupled with the much higher specific energy consumption enables the designers to tolerate orders of magnitude higher cost with the ultimate objective of achieving savings in weight [1,2]. Considering a much higher payoff over a system’s life cycle, the tolerable material prices in the aerospace sector is about three to five orders of magnitude higher than those in the automobile industry [3]. The much higher payoff for reduction in weight of aircraft, and especially the aircraft structures, can be attributed to the much lower payload capability when compared one-on-one with the land-based vehicles. For example, a Boeing 747 commercial aircraft carries about 100 tons of fuel, which is approximately one-third the
∗ Corresponding author. Tel.: +1 330 972 6196; fax: +1 330 972 6027. E-mail address: [email protected] (T.S. Srivatsan). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.06.086
take-off weight of the jumbo-jet. If the consumption of fuel is lowered by only 10 percent, it is then possible to increase the pay-load of the Boeing 747 airliner by about ten times [3]. Compared to the family of steels and aluminum alloys, the alloys of titanium can be safely considered to be a much younger structural material. The first alloys of titanium were developed at the end of 1940s in USA. Among these was the classic titanium alloy, Ti–6Al–4V, which has grown in strength to capture a large portion of the applications relevant to the aerospace sector [4]. This alloy has over the years since its inception and based on its selection and use for a plethora of applications grown in stature to be referred to in industry circles as the ‘work-horse’ alloy. The development of alloys based on the light metal titanium was largely stimulated by the aerospace industry. The outstanding properties of the alloys of titanium, most notably and specifically Ti–6Al–4V, include a combination of (i) high specific strength (/), (ii) good fracture toughness, and (iii) a balance of strength, ductility and fatigue crack growth resistance [5,6]. Further, it possesses other excellent properties, such as, weldability, excellent resistance to corrosion, and good oxidation resistance [7]. Some of the titanium alloys exhibit a room temperature tensile strength of more than 1500 MPa and can be safely used at temperatures as high as 600 ◦ C. Consequently, the alloys of titanium, specifically Ti–6Al–4V, are found applicable for a spectrum of aerospace application where a combination of weight, strength, and corrosion resistance are required and the high temperature stability of aluminum
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Table 1 Nominal chemical composition of the Ti–6Al–4V alloy (in weight percent). Material
Ti
Al
N
V
C
Fe
H
O
Ti–6Al–4V
90.0
6.0
0.05
4.0
0.1
0.4
0.02
0.20
alloys, high strength steels and the nickel-base super-alloys are insufficient [8]. Since its development and emergence around the late 1940s, the Ti–6Al–4V alloy has grown in strength and preferential selection or choice for use in industries spanning aircraft structural and engine components, biomedical to include surgical implants, and petro-chemical plants due essentially to its excellent combination of mechanical properties, fatigue crack growth resistance and even oxidation resistance. Since its emergence the fatigue behavior of Ti–6Al–4V, particularly fatigue crack growth, has been studied by a number of investigators [9–23]. These studies have essentially shown that (a) microstructure of the alloy to include crystallography and texture [24–27], (b) stress ratio [28–30], and (c) environment [13,31,32] exert a profound influence on fatigue response. For the Ti–6Al–4V alloy categorized as having an alpha plus beta (␣ + ) dual phase, the microstructure is known to exert a dominating influence on fatigue properties. The alpha (␣) phase consists of a hexagonal close-packed (HCP) crystal structure, while the beta () phase is essentially body-centered cubic (BCC). It is generally believed and supported by sufficient experimental data that refining the microstructure helps in an improvement in the fatigue strength of the smooth specimens [9–13]. However, the overall resistance to fatigue crack propagation was found to decrease [14,15]. The dual phase [i.e., alpha plus beta] Ti–6Al–4V alloy when properly heat treated has to offer acceptable combination of strength and ductility when properly heated treated, which makes it stronger than the near-alpha phase alloy and near-beta phase alloy counterpart and a preferred candidate for fatiguecritical applications. In this paper, the results of a recent study aimed at understanding the high-cycle fatigue response and final fracture behavior of samples taken from a welded plate of Ti–6Al–4V alloy are presented and discussed. The as-welded microstructure of the alloy in the region immediately adjacent to the weld was characterized for: (i) the nature and morphology of phases present, and (ii) the presence and distribution of other intrinsic microstructural features. The high-cycle fatigue response of the alloy was determined at a load ratio of 0.1 over a range of maximum stress ( max ) and resultant fatigue life (Nf ). The high-cycle fatigue properties and final fracture behavior of the alloy are discussed in light of the conjoint and mutually interactive influences of maximum applied stress, test specimen orientation, intrinsic microstructural features, and stress-deformation-microstructural feature interactions. Of particular interest in this research study was to systematically evaluate and rationalize the role of test specimen orientation (longitudinal versus transverse) in governing damage initiation, damage propagation and overall cyclic resistance of the alloy microstructure quantified in terms of cyclic fatigue life and final fracture behavior.
ufacturer in the fully annealed condition. The nominal chemical composition of the alloy is provided in Table 1. The tensile and fatigue test specimens were precision machined from both the longitudinal and transverse orientations of the two annealed plates that were welded together using the technique of butt welding. The longitudinal specimens were machined with the major stress axis parallel to the rolling direction of the aswelded titanium alloy plate, while the transverse specimens were machined with the major stress axis perpendicular to the rolling direction of the as-welded titanium alloy plate. The required sections for machining the test specimens were taken from a region of the plate immediately adjacent to the weld (Fig. 1). The test specimens were prepared in conformance with standards outlined in the standard ASTM E-8. At the gage section, the test specimens measured 3.125 mm in diameter and 12.5 mm in length. Final surface preparation was achieved by mechanically polishing the gage section of the test specimens with progressively finer grades of silicon carbide embedded emery paper (320-grit, 400-grit and 600-grit), so as to minimize the effects and/or contributions from surface irregularities, such as, scratches, residual machine marks, and surface finish. 3. Experimental procedures 3.1. Initial microstructure characterization An initial characterization of the microstructure of the aswelded Ti–6Al–4V alloy was done using a low magnification optical microscope. Samples were cut from the welded plate of the Ti–6Al–4V alloy and mounted in bakelite. The mounted samples were then wet ground on progressively finer grades of silicon carbide impregnated emery paper using copious amounts of water both as a lubricant and as a coolant. Subsequently, the ground samples were mechanically polished using five-micron diamond solution. Fine polishing to a perfect mirror-like finish of the surface of each titanium specimen was achieved using one-micron diamond solution as the lubricant. The polished samples were then etched using a reagent that is a solution mixture of 5-ml of nitric acid (HNO3 ), 10 ml of hydrofluoric acid (HF) and 85 ml of water (H2 O). The polished and etched surface of the samples of Ti–6Al–4V were observed in an optical microscope and photographed using a bright field illumination technique.
2. Materials and sample preparation The material chosen for this research study was the widely preferred and chosen alloy Ti–6Al–4V. The alloy was provided by Allegheny Technologies ATI Wah Chang (based in Oregon, USA). Owing to the relative ease of production coupled with good hot workability and receptiveness to heat treatment, the Ti–6Al–4V alloy can be easily solution heat treated and annealed to achieve the desired strength properties. The alloy was provided by the man-
Fig. 1. A schematic of the longitudinal and transverse flat test specimen used for Mechanical testing (tensile and cyclic fatigue).
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Fig. 2. Optical micrographs showing the key micro-constituents, their morphology and distribution, in the Ti–6Al–4V alloy at three different magnifications.
3.2. Tensile and cyclic fatigue tests The tensile properties were determined by uniaxial tensile tests performed on a fully automated, closed-loop servo-hydraulic mechanical test machine [INSTRON-8500 Plus] using a 100 KN load cell. The tests were conducted at room temperature (25 ◦ C) and in the laboratory air environment (Relative Humidity of 55 pct.). The test specimens were deformed in uniaxial tension at a constant strain rate of 0.0001/s. An axial 12.5-mm gage length clip-on type extensometer was attached to the test specimen at the gage section using rubber bands. The stress and strain measurements, parallel to the load line, and the resultant mechanical properties, namely, (i) stiffness, (ii) yield strength, (iii) ultimate tensile strength, (iv) failure stress, and (v) ductility (strain-to-failure), was provided
Fig. 3. Optical micrographs showing minimal influence in microstructure of the heat affected zone and the base metal.
as a computer output by the controller unit of the mechanical test machine. The stress-amplitude controlled high-cycle fatigue tests were performed using a sinusoidal waveform at a stress ratio [R = minimum stress ( min )/maximum stress ( max )] of 0.1. The stress-controlled fatigue tests were conducted at a frequency of 5 Hz. At the chosen stress ratio [R = min / max ], the fatigue tests were conducted over a range of stress amplitudes to establish the variation of maximum stress ( max ) with cyclic fatigue life (Nf ). For the purpose of this investigation the fatigue limit was taken to be the stress below which fatigue failure does not occur even after 106 (one-million) cycles.
3.3. Characterization of damage and analysis of failure The fracture surfaces of the cyclically deformed and failed test specimens of the Ti–6Al–4V alloy were thoroughly examined in a scanning electron microscope (SEM) to determine the macroscopic fracture mode and to concurrently characterize the fine scale topography of the cyclic fatigue fracture surface in order to establish both the macroscopic mechanisms and concurrent microscopic mechanisms governing fracture. The distinction between the macroscopic mode of failure and microscopic fracture mechanisms is based entirely on the magnification level at which the observations are made. The macroscopic mode refers to the overall nature of the failure while the microscopic mechanisms relate to the local failure processes, such as: (i) microscopic and macroscopic void for-
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27 25 – –
Tensile ductility ln(Ao /Af ) (%)
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mation, (ii) growth of the microscopic and macroscopic voids and their eventual coalescence, and (iii) nature, intensity and severity of the fine microscopic cracks and macroscopic cracks dispersed throughout the fracture surface. The samples for observation in the scanning electron microscope were obtained from the failed fatigue specimens by sectioning parallel to the fracture surface.
Reduction in area (%)
4. Results and discussion
148 167 140–158 130–180 921 1029 875–995 800–1100
1020 1153 965–1090 900–1200
ksi MPa
134 149 127–144 115–160
ksi
126 137 105 114
GPa msi
Carpenter technology corporation ASM handbook
Longitudinal Transverse – – Ti–6Al–4V
18 20 15 17
Yield strength Elastic modulus Orientation Material
Table 2 Room temperature tensile properties of Ti–6Al–4V [results are mean values based on duplicate tests].
UTS
MPa
7.8 11.5 16 –
Elongation GL = 0.5 (%)
23.8 21.7 46 –
4.1. Initial microstructure characterization The microstructure of the material under consideration is one of the major factors in determining the mechanical properties like yield strength, tensile strength, ductility, fatigue resistance and resultant fracture behavior. The optical microstructure of the welded Ti–6Al–4V alloy plate is shown in Fig. 1 at three different magnifications. A careful observation of the microstructure over a range of low magnifications in the optical microscope revealed a two-phase duplex microstructure consisting of the HCP equiaxed primary alpha (␣) and interspersed grains containing HCP secondary ␣ and BCC  arranged in a lamellar structure. The primary near equiaxed shaped alpha (␣) grains (light in color) were well distributed in the lamellar matrix of secondary ␣ and transformed beta (dark in color) (Fig. 2). The size of both the alpha (␣) grains and the interspersed grains containing secondary ␣ and  range in size from 10 m to 20 m. The presence of trace amounts of aluminum and oxygen in the Ti–6Al–4V alloy contributes to stabilizing and strengthening of the HCP alpha (␣) phase. This is particularly beneficial for enhancing the hardenability, increasing strength, and improving the response kinetics of the alloy to heat treatment [8]. In Fig. 3 is shown the microstructure of the as-welded alloy plates. 4.2. Tensile properties The room temperature tensile properties of the Ti–6Al–4V are summarized in Table 2. The results are the mean values based on duplicate tests. The elastic modulus, yield strength, ultimate tensile strength, elongation to failure and strength at failure (fracture) were provided as an output of the computer-based data acquisition and test control system. The yield strength was determined by identifying the stress at a point on the engineering stress versus engineering strain curve where a straight line drawn parallel to the elastic portion of the curve at 0.2% offset intersects the curve. The ductility is reported as elongation-to-failure over a gage length of 12.5 mm (0.5 in.). This elongation was measured using a clipon extensometer that was attached to the gage section of the test specimen. The elastic modulus of the test specimens machined immediately adjacent to the weld region of the alloy plate was 126 GPa in the longitudinal (L) orientation and 136 GPa in the transverse (T) orientation. The yield strength was 921 MPa (134 ksi) in the longitudinal (L) orientation, which is approximately 12 percent lower than in the transverse (T) orientation [1029 MPa (149 ksi)]. The ultimate tensile strength is 1020 MPa (148 ksi) in the longitudinal (L) orientation, which is 12 percent lower than in the transverse (T) orientation [1153 MPa (167 ksi)]. The ultimate tensile strength of this material is only marginally higher than the yield strength indicating the occurrence of strain hardening or work hardening beyond yield. Ductility quantified in terms of elongation-to-failure was 7.8% in the longitudinal (L) orientation and 11.5% in the transverse (T) orientation. The reduction-in-area (RA) experienced by the test specimen was as high as 24% in the longitudinal (L) orientation and 22% in the transverse orientation. The yield strength and tensile strength values of the alloy conform well with the val-
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Fig. 4. Variation of maximum stress ( max ) with fatigue life (Nf ) for Ti–6Al–4V.
ues obtained and reported by the manufacturer (ATI Wah Chang) and recorded in ASM Metals Handbook [33]. 4.3. Variation of cyclic stress with fatigue life: high-cycle fatigue properties A gradual degradation in the cyclic strength of a material through the accumulation of damage at the ‘local’ level is known as fatigue. The fatigue life (Nf ) of a structure, or component, is defined as the sum of the number of cycles to crack initiation (Ni ) plus the number of cycles to crack propagation (Np ), which eventually culminates in failure by fracture. A comprehensive study of the cyclic fatigue response of the welded Ti–6Al–4V alloy plates is important for its selection and subsequent use in an array of performancecritical applications. Thus, a goal of this investigation was to extend prior research on high-cycle fatigue of non-fabricated material into an area involving a fabricated structural component in an effort to reflect upon the conditions while in actual use or service. A comprehensive understanding of the fatigue behavior and related damage mechanisms governing failure is rendered difficult by the underlying intrinsic microstructural features and their specific role in governing cyclic deformation and fracture processes during cyclic loading. It has been established in earlier studies that the fatigue properties of the commercially available titanium alloys are governed by the conjoint influences of (a) surface defects, (b) crystallographic texture, and (c) intrinsic microstructural effects to include the following: (i) size and morphology of the interspersed grains with the ␣ and  phases arranged in a lamellar structure, (ii) size of the alpha platelets, and (iii) the presence and role of other intrinsic microstructural features [34–40]. Alloy microstructures that favor a low resistance to crack initiation, even at the low values of applied stress, will usually yield the best endurance during high-cycle fatigue testing and even during tensile resistance. The cyclic fatigue test is the most widely used technique to establish the endurance limit of a chosen metal by essentially determining the variation of maximum stress ( max ) or stress amplitude (/2), with fatigue life as quantified by the number of cycles to failure (Nf ). Metals, alloys and other materials based on metal matrices, which are used at stress levels below the endurance limit can be cycled indefinitely [17]. The endurance limit for this non-ferrous metal is taken to be 106 cycles. The high-cycle fatigue test is the most widely used technique to establish the endurance limit of the chosen metal by essentially determining the variation of maximum stress ( max ) with fatigue life (Nf ). For small, highly stressed components, fatigue life is often controlled by cyclic plasticity at the ‘local’ level coupled with the relatively rapid growth of the fine microscopic cracks through the microstructure of the alloy. For larger components that operate at low stress levels, the fatigue life is initially controlled by micro-
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Fig. 5. Variation of maximum elastic strain ( max /E) with fatigue life (Nf ) for Ti–6Al–4V.
scopic crack formation and subsequently by their growth through the alloy microstructure and eventual coalescence to form one or more macroscopic cracks. The kinetics of crack nucleation in this Ti–6Al–4V alloy is governed by the mutually interactive influences of the following: (i) the rate of cooling, (ii) volume fraction and size of the two key micro-constituent phases, i.e., alpha and beta, and (iii) local stress concentration effects or surface residual stress [17–19]. At the load ratio studied (R = 0.1), the variation of maximum stress ( max ) with cyclic fatigue life (Nf ) reveals increasing fatigue life with a decrease in stress amplitude. This trend is shown by most non-ferrous metals and their composite counterparts. At equivalent maximum stress, the fatigue life of the transverse (T) specimen is marginally greater than the longitudinal (L) counterpart (Fig. 4). In an attempt to understand and rationalize the influence of specimen ductility on high-cycle fatigue response the test data is re-plotted as the variation of maximum elastic strain ( max /E) as a function of cycles-to-failure (Nf ). The maximum elastic strain is taken to be the ratio of maximum stress ( max ) to the elastic modulus (E) of Ti–6Al–4V alloy in that specific orientation. This is shown in Fig. 5. At equivalent values of maximum elastic strain particularly at low values of strain and resultant enhanced fatigue life the longitudinal specimen revealed a marginal improvement in fatigue life over the transverse counterpart (Fig. 6). In an attempt to understand the specific role of alloy microstructure in high-cycle fatigue life at the chosen load ratio of 0.1, the test data is plotted to take into consideration the strength of the Ti–6Al–4V alloy plate in the as-welded condition. The maximum stress is normalized by the yield strength of the alloy plate and reveals minimal difference in fatigue life of the two orientations,
Fig. 6. Variation of the ratio of maximum stress/yield stress ( max / y ) to fatigue life (Nf ) for Ti–6Al–4V.
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Fig. 7. Variation of the ratio of maximum stress/ultimate tensile stress ( max / UTS ) to fatigue life (Nf ) for Ti–6Al–4V.
i.e., longitudinal and transverse of the welded alloy plate, within the limits of experimental scatter. Variation of maximum stress normalized by the ultimate tensile strength of the alloy plate in that specific orientation with fatigue life (Nf ) reveals that at near equivalent values of the ratio of max / UTS , the influence of test
Fig. 8. Comparing the variation of maximum stress ( max ) with fatigue life (Nf ) for Ti–6Al–4V alloy test specimens and ASM Handbook for a given stress ratio of (R = 0.1).
specimen orientation on cyclic fatigue life is minimal at all values of the ratio (Fig. 7). Comparing the fatigue test results obtained in this research study with the stress versus fatigue life response of the alloy documented in the published literature is shown in Fig. 8. This figure reveals the high-cycle fatigue resistance of both the longitudinal
Fig. 9. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Longitudinal) deformed in cyclic fatigue at a maximum stress of 829 MPa and resultant fatigue life (Nf ) of 255,334 cycles, showing: (a) overall morphology of failure. (b) The region of crack initiation and radial outward progression of fatigue damage. (c) The transgranular surface: flat and near featureless in the region of early crack growth. (d) Shallow dimples covering the region of overload reminiscent of locally ductile failure mechanism.
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Fig. 10. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Longitudinal) deformed in cyclic fatigue at a maximum stress of 811 MPa and resultant fatigue life (Nf ) of 674,128 cycles, showing: (a) overall morphology of failure. (b) The region between stable and unstable crack growth showing a healthy population of fine microscopic voids and dimples. (c) High magnification of (b) showing a population of macroscopic and fine microscopic voids. (d) Void coalescence to form a fine microscopic crack and shallow dimples indicative of locally ductile failure mechanisms.
and transverse test specimens taken from the as-welded plate to be noticeably better than the high-cycle fatigue resistance of the as-prepared and annealed Ti–6Al–4V alloy. At equivalent values of maximum stress the improvement was of the order of 100–200 percent. This noticeable difference in cyclic fatigue resistance is ascribed to intrinsic microstructural differences between the conventional alloy and the material chosen for this research study and its contribution to governing cyclic fatigue resistance. 4.4. An analysis and understanding of damage initiation, damage propagation and cyclic fracture The fracture surfaces of the high-cycle fatigue test specimens were carefully examined in a JEOL scanning electron microscope (SEM) at the following levels of magnification: (a) Low magnification in an attempt to identify the nature of microscopic crack initiation, the extent of early crack growth through the alloy microstructure and the kinetics of stable crack growth. (b) Gradually increasing levels of magnification for the purpose of studying the nature of damage initiation, nature and depth of stable crack propagation, and other fine scale features on the fracture surface in the region of overload.
For the Ti–6Al–4V material, the fatigue fracture surfaces revealed only a marginal difference in the topographies at the different levels of maximum stress and resultant fatigue life for both the longitudinal and transverse specimens. On a microscopic scale, the nature and morphology of the fracture surfaces did not vary extensively with maximum stress and fatigue life. Only representative fractographs of the fatigue fracture surfaces of samples of Ti–6Al–4V taken from the welded plate, for both the longitudinal and transverse specimens, at different values of maximum stress and corresponding cyclic fatigue life, are shown in Figs. 9–14.
4.4.1. Longitudinal orientation At a maximum stress of 829 MPa, the longitudinal specimen had a fatigue life of 255,334 cycles. Low magnification observations in the scanning electron microscope revealed the existence of well defined and distinct region of crack initiation and a short region representative of stable crack propagation (Fig. 9a). Early microscopic crack growth was observed to be extending radially outward from the point marking the onset of crack initiation (Fig. 9b). A careful examination of the region of stable crack propagation at the higher allowable magnifications of the SEM revealed the region to be flat and featureless, containing a population of fine microscopic cracks that were randomly distributed through the region (Fig. 9c).
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Fig. 11. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Longitudinal) deformed in cyclic fatigue at a maximum stress of 811 MPa and resultant fatigue life (Nf ) of 674,128 cycles, showing: (a) The transgranular surface in the region of early crack growth. (b) High magnification of (a) showing pockets of fine shallow striations indicative of localized micro-plastic deformation. (c) Well defined striations in the region of stable crack growth indicative of micro-plastic deformation.
For almost all the samples cyclically deformed at varying levels of maximum stress, the overload region revealed the presence of a population of shallow dimples indicative of locally ductile failure mechanisms (Fig. 9d). For the test specimen that was cyclically deformed at a maximum stress of 811MPa the resultant fatigue life was 674,128 cycles. The fracture surfaces of this test specimen revealed distinctly the regions of crack initiation and overload (Fig. 10a). In the terminal stages of the region of stable crack growth through the alloy microstructure, a transition region between stable and unstable crack growth was distinctly evident (Fig. 10b). High magnification observations of the transition region revealed a population of macroscopic voids and fine microscopic voids (Fig. 10c). Continued cyclic loading led to a progressive growth of the fine microscopic voids and their eventual coalescence resulting in the formation of one or more fine microscopic cracks. Half of these microscopic voids are the shallow dimples visible on the fracture surface (Fig. 10d). The transgranular region (Fig. 11a), when observed at high magnifications, revealed (a) pockets of fine shallow striations in the region of early microscopic crack growth, and (b) well defined striations (Fig. 11b). This feature is indicative of the localized micro-plastic deformation mechanisms in the region of stable crack growth. Thus, for the longitudinal orientation, the microscopic fracture characteristics observed for the test sample that was cyclically
deformed at high maximum stress, resultant low fatigue life, and for the test sample that was cyclically deformed at a low maximum stress and resultant enhanced fatigue life, are essentially similar. 4.4.2. Transverse orientation Following the same trend as the specimens oriented in the longitudinal direction, the test specimens prepared from the transverse orientation of the as welded Ti–6Al–4V alloy plate exhibited almost identical fatigue lives at the higher level of maximum stress. At a maximum stress of 962 MPa and resultant fatigue life of 40,664 cycles, the fracture surface features are shown in Fig. 12. The overall morphology comprised of (a) distinct regions of crack initiation and growth, (b) a short region of stable crack growth, and (c) overload (Fig. 12a). The region of early crack growth revealed radial propagation of the fatigue damage from the onset of initiation (Fig. 12b). High magnification observations in the region of stable crack growth revealed fine and shallow striations indicative of localized micro-plastic deformation (Fig. 12c). The region of overload revealed a population of both fine microscopic voids and macroscopic voids. For a specimen subjected to a maximum stress of 900 MPa and resultant enhanced fatigue life of 425,984 cycles, the overall fracture surface morphology revealed distinctive regions of crack initiation and radial propagation of the cracks across the transgranular surface region (Fig. 13a). High magnifica-
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Fig. 12. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Transverse) deformed in cyclic fatigue at a maximum stress of 962 MPa and resultant fatigue life (Nf ) of 40,664 cycles, showing: (a) overall morphology of failure. (b) The region of early crack growth: flat and near featureless. (c) Pockets of shallow striations in the region of stable crack growth. (d) A population of fine microscopic and macroscopic surface voids in the region of overload.
tion observations in the region of early crack propagation (Fig. 13b) revealed a random distribution of fine microscopic cracks (Fig. 13c). The characteristic features on the overload surface were fine microscopic voids and a population of shallow dimples, and these were observed in test specimens that were cyclically deformed at both high maximum stress and low maximum stress (Figs. 12d and 13d). Repeated pockets of very fine and shallow striations were found at random locations in the region of stable crack growth indicative of the highly localized ductile failure mechanisms (Fig. 14). 4.5. The mechanisms governing the initiation of damage, propagation of damage and fracture at microscopic level It is well known that under cyclic stress control, the resultant fatigue life is dependent on (a) the nature of crack initiation, and (b) the nature and severity of crack propagation through the alloy microstructure. Much of the fatigue life is spent or utilized for the initiation of fine microscopic cracks. The intrinsic microstructural features that were formed during initial heat treatment and during subsequent fabrication of the welded alloy plate do exert an influence on cyclic fatigue response and resultant fatigue life of the candidate alloy. The microstructure of the Ti–6Al–4V alloy material chosen for this study is taken from a region, which can be characterized as having a duplex microstructure consisting of near equiaxed alpha (␣) and transformed beta () phases. In order to
account for requirements of continuity, slip in the HCP primary ␣ phase accounts for majority of the plastic deformation in this alloy. The lamellar structure merely acts to inhibit the slip in the primary ␣ phase and this occurs at the ␣- phase boundaries. An earlier study on ␣- alloy titanium alloy has shown that slip initiates in the primary ␣ phase, which is mechanically softer than the  lamellar phase. Further, the mechanical properties of this and other two phase alloys of titanium are very sensitive to the geometrical arrangement of the phases and the crystallographic texture. Other companion and subsequent studies by researchers have shown that fatigue cracking in the Ti–6Al–4V alloy initiates both at and along the grain boundaries. This was attributed to the presence of a large area fraction of weakened high angle grain boundaries parallel to the longitudinal direction. The preference for crack initiation and early crack propagation along the grain boundaries results in a progressive loss of the through thickness constraint causing an essentially plane strain fracture process to be divided into several plane stress fractures. During continued cyclic deformation of the alloy specimens, a gradual degradation of the intrinsic microstructural features results in the formation of “localized” failure sites coupled with the accumulation of strain at the microscopic level. When the strain accumulation and resultant stress concentration at the localized level exceeds a critical value, microscopic crack initiation is favored to occur, which is exacerbated by the continued application of stress culminating
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Fig. 13. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Transverse) deformed in cyclic fatigue at a maximum stress of 900 MPa and resultant fatigue life (Nf ) of 425,984 cycles, showing: (a) overall morphology of failure showing region of initiation and radial progression of fatigue damage. (b) High magnification of region of early crack growth. (c) High magnification observation of the region of early crack growth showing a random distribution of fine microscopic cracks. (d) A population of voids and dimples in the region of overload.
in the progressive accumulation of micro-plastic deformation at the ‘local’ level. Continuous cyclic deformation ably aided by the continuous propagation of the fine microscopic cracks resulting from the growth and eventual coalescence of the microscopic voids leads to the formation of macroscopic cracks. A gradual growth or propagation of both the fine microscopic and macroscopic cracks through the alloy microstructure results in a progressive drop in load-carrying capability of the test specimen culminating in failure. Based on a careful and comprehensive observation of the fracture surfaces of the test specimens taken from both the longitudinal and transverse orientations of the as-welded titanium alloy plates and cyclically deformed at specific values of maximum stress, the fracture plane of the cyclically deformed and failed test specimens was essentially normal to the stress or loading axis, suggesting the importance of tensile stress in promoting fracture. Few of the microscopic voids that were created by the presence of intrinsic microstructural features did not grow extensively in the direction of applied stress, which is generally the case for ductile failure of metallic materials. The lack of extensive growth of the fine microscopic voids suggests that the overall strain that induces fracture at both the microscopic and macroscopic level is controlled by (i) the strain required for void nucleation, (ii) crack propagation strain, and (iii) the eventual strain required for their linkage.
Fig. 14. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Transverse) deformed in cyclic fatigue at a maximum stress of 823 MPa and resultant fatigue life (Nf ) of 581,951 cycles, showing pockets of shallow striations reminiscent of local micro-plastic deformation.
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5. Conclusions A study of the high-cycle fatigue and final fracture behavior of test specimens machined from a Ti–6Al–4V as welded plate for the two different orientations, i.e. longitudinal and transverse provides the following key observations: 1. The Ti–6Al–4V alloy exhibited a duplex microstructure consisting of the near equiaxed alpha (␣) and the interspersed grains of ␣ and  phases. The primary near equiaxed shaped alpha (␣) grains was well distributed in a lamellar matrix with transformed beta. 2. The yield strength and ultimate tensile strength of Ti–6Al–4V is nearly equal both in the longitudinal orientation and in the transverse orientation indicating a more or less isotropic nature of the material. In both the longitudinal and transverse orientations, the tensile strength of the as-welded alloy plate is higher than the yield strength indicating the occurrence of work hardening beyond yield. 3. There is only a marginal influence of microstructure on highcycle fatigue life of the as-welded alloy plate in the two orientations. The variation of maximum stress with cyclic fatigue life (Nf ) revealed a trend shown by most non-ferrous metallic materials. At equivalent values of maximum stress, at a load ratio of 0.1, the fatigue life of the longitudinal and transverse specimens show minimal difference at all levels of stress, with the transverse specimen exhibiting marginally improved fatigue life when compared to the longitudinal counterpart. 4. At the chosen load ratio of 0.1 the fatigue fracture surfaces revealed minimal to no difference in the nature, distribution and volume fraction of the intrinsic features on the fracture surface as a function of maximum stress and resultant fatigue life. The region of crack initiation and early microscopic crack growth and a short, yet distinct, region of stable crack growth were essentially flat and transgranular. A population of fine microscopic and macroscopic cracks was evident in the region of stable crack growth approaching overload. Also evident in this region were pockets of well defined striations indicative of the occurrence of micro-plastic deformation at the local level. The region of overload was covered with a population of shallow dimples, fine microscopic voids and isolated macroscopic void and cracks, features reminiscent of the locally governing failure mechanisms. During repeated cyclic loading, the fine microscopic cracks coalesce to form macroscopic crack. References [1] M. Peters, R. Braun, 4th European Conference on Advanced Materials and Processes – EUROMAT 95, Associazione Italiana di Metallurgia, Milano, 1995, p. 55.
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[2] M. Peters, Titanium World 2 (1995) 15. [3] M. Peters, J. Kumpfert, C.H. Ward, C. Leyens, Advanced Engineering Materials 5 (6) (2003) 419–427. [4] D. Eylon, S.R. Seagle, Titanium 99: Science and Technology, CRISM Prometey, St Petersburg, 2000, pp. 37–41. [5] G. Thomas, V. Ramachadra, K.V. Nagarajan, B. Pant, B.K. Sarkar, R. Vasudevan, Welding Research Supplement 8 (1989) 336s. [6] J.L. Gilbert, H.R. Piehler, Metallurgical Transactions 20A (1989) 1715–1720. [7] D. Helm, in: R.R. Boyer, D. Eylon, G. Lutjering (Eds.), Fatigue Behavior of Titanium Alloys, TMS: The Minerals, Metals and Materials Society, Warrendale, Pennsylvania, USA, 1998, pp. 291–298. [8] J.M. Donachie, Titanium: A Technical Guide, ASM International, Metals Park, Ohio, USA, 1988. [9] J.J. Lucas, P.P. Konieczny, Metallurgical Transactions 2 (1971) 911. [10] A.W. Bowen, C.A. Stubbington, Titanium Science and Technology, vol. 3, Plenum Press, New York, NY, 1973, p. 2097. [11] LucasF J.J., Titanium Science and Technology, vol. 3, Plenum Press, New York, 1973, p. 2081. [12] C.A. Stubbington, A.W. Bowen, Journal of Materials Science 9 (1974) 941. [13] M. Peters, A. Gysler, G. Lutjering, Titanium 80: Science and Technology, vol. 3, TMS–AIME, Warrendale, PA, USA, 1980, p. 1777. [14] H. Margolin, J.C. Williams, J.C. Chestnut, G. Lutjering, Titanium 80: Science and Technology, vol. 1, TMS–AIME, Warrendale, PA, USA, 1980, p. 169. [15] G.R. Yoder, L.A. Cooley, T.W. Crooker, Engineering Fracture Mechanics 11 (1979) 805. [16] J. Oh, N.J. Kim, S. Lee, E.W. Lee, Materials Science and Engineering A340 (2003) 232–242. [17] L.W. Tsay, C.Y. Tsay, International Journal of Fatigue 19 (10) (1997) 713–718. [18] R.J. Morrissey, D.L. McDowell, T. Nicholas, International Journal of Fatigue 21 (1999) 679–685. [19] R.J. Morrissey, D.L. McDowell, T. Nicholas, International Journal of Fatigue 23 (2001) S55–S64. [20] K. Tokaji, Scripta Materialia 54 (2006) 2143–2148. [21] I. Bantounas, T.C. Lindley, D. Rugg, D. Dye, Acta Materialia 55 (2007) 5655–5665. [22] A. Yuen, S.W. Hopkins, G.R. Leverat, C.A. Rau, Metallurgical Transactions 5 (8) (1974) 1833–1842. [23] K.S. Ravichandran, E.S. Dwarakadasa, Scripta Metallurgica 23 (10) (1989) 1685–1690. [24] M. Peters, A. Gysler, G. Lutjering, Metallurgical Transactions 15A (1984) 1597–1605. [25] H. Margolin, J.C. Chestnutt, G. Lutjering, J.C. Williams, Titanium 80, Science and Technology, TMS–AIME, Warrendale, PA, USA, 1980, p. 169. [26] J.C. Chestnutt, A.W. Thompson, J.C. Williams, Titanium 80: Science and Technology, TMS–AIME, Warrendale, PA, USA, 1980, p. 1875. [27] P.E. Irving, C.J. Beevers, Materials Science and Engineering 14 (1974) 229. [28] G.T. Gray, G. Lutjering, Titanium Science and Technology 4 (1985) 2251. [29] A.R. Rosenfield, Engineering Fracture Mechanics 9 (1977) 510. [30] S. Dubey, A.B.O. Soboyejo, W.O. Soboyejo, Acta Materialia 45 (7) (1997) 2777–2787. [31] R.J.H. Wanhill, The Aeronautical Journal of the Royal Aeronautical Society (1977) 68. [32] P.E. Irving, C.J. Beevers, Metallurgical Transactions 5 (1974) 391–400. [33] Titanium Alloys, Materials Properties Handbook, American Society for Materials International, Materials Park, Ohio, USA, 1994. [34] A.W. Bowen, Scripta Metallurgica 11 (1977) 17. [35] A.W. Bowen, Titanium Science and Technology, vol. 2, Plenum Press, New York, NY, 1973, p. 1271. [36] R.J.H. Wanhill, Metallurgical Transactions 7A (1976) 1365. [37] M.H. Muller, H.W. Knott, Review Science Instruments 25 (1954) 1115. [38] G. Lutjering, M. Peters, R.I. Jafee, Mechanical Properties of a Titanium Blading. Alloy, EPRI-CS-2933, Electric Power Research Institute, Palo Alto, CA, 1983. [39] M. Peters, G. Lutjering, G. Ziegler, Zeitschrift f Metallkunde 74 (1983) 274. [40] G. Welsch, G. Lutjering, K. Gazioglu, W. Bunk, Metallurgical Transactions 8A (1977) 169.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
A study of cyclic fatigue, damage initiation, damage propagation, and fracture of welded titanium alloy plate T.S. Srivatsan a,∗ , U. Bathini b , A. Patnaik b , T. Quick c a b c
Department of Mechanical Engineering, The University of Akron, Akron, OH 44325-9303, USA Department of Civil Engineering, The University of Akron, Akron, OH 44325-9303, USA Department of Geology, The University of Akron, Akron, OH 44325, USA
a r t i c l e
i n f o
Article history: Received 19 May 2010 Accepted 28 June 2010
Keywords: Welded titanium alloy Damage initiation Damage propagation Fatigue life Microstructure Fracture
a b s t r a c t In this paper, the influence of test specimen orientation and microstructure on cyclic stress-amplitude controlled fatigue response, damage initiation, damage propagation and fracture behavior of samples taken from a welded plate of titanium alloy is presented and discussed. Test specimens from the chosen alloy were prepared from an as-welded plate of the material with the stress axis both parallel (longitudinal) and perpendicular (transverse) to the deformed (rolling) direction of the plate. The test specimens were cyclically deformed at different values of maximum stress at a constant load ratio of 0.1, and the resultant cycles-to-failure was recorded. The fracture surfaces of the deformed and failed test specimens were examined in a scanning electron microscope to establish the macroscopic fracture mode, the intrinsic features on the fatigue fracture surface and the role of applied stress-microstructural feature interactions in establishing the microscopic mechanisms governing failure. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The ever increasing scarcity of resources coupled with their growing expense demands both attention and action towards a reduction in the consumption of energy when dealing with the transportation of passengers and goods. It is here that the aerospace sector comes into to picture in playing an important and key role with specific respect to the development, emergence, application and commercialization of new and improved materials. In direct comparison with the land-based transportation systems, the much lower system quantities coupled with the much higher specific energy consumption enables the designers to tolerate orders of magnitude higher cost with the ultimate objective of achieving savings in weight [1,2]. Considering a much higher payoff over a system’s life cycle, the tolerable material prices in the aerospace sector is about three to five orders of magnitude higher than those in the automobile industry [3]. The much higher payoff for reduction in weight of aircraft, and especially the aircraft structures, can be attributed to the much lower payload capability when compared one-on-one with the land-based vehicles. For example, a Boeing 747 commercial aircraft carries about 100 tons of fuel, which is approximately one-third the
∗ Corresponding author. Tel.: +1 330 972 6196; fax: +1 330 972 6027. E-mail address: [email protected] (T.S. Srivatsan). 0921-5093/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2010.06.086
take-off weight of the jumbo-jet. If the consumption of fuel is lowered by only 10 percent, it is then possible to increase the pay-load of the Boeing 747 airliner by about ten times [3]. Compared to the family of steels and aluminum alloys, the alloys of titanium can be safely considered to be a much younger structural material. The first alloys of titanium were developed at the end of 1940s in USA. Among these was the classic titanium alloy, Ti–6Al–4V, which has grown in strength to capture a large portion of the applications relevant to the aerospace sector [4]. This alloy has over the years since its inception and based on its selection and use for a plethora of applications grown in stature to be referred to in industry circles as the ‘work-horse’ alloy. The development of alloys based on the light metal titanium was largely stimulated by the aerospace industry. The outstanding properties of the alloys of titanium, most notably and specifically Ti–6Al–4V, include a combination of (i) high specific strength (/), (ii) good fracture toughness, and (iii) a balance of strength, ductility and fatigue crack growth resistance [5,6]. Further, it possesses other excellent properties, such as, weldability, excellent resistance to corrosion, and good oxidation resistance [7]. Some of the titanium alloys exhibit a room temperature tensile strength of more than 1500 MPa and can be safely used at temperatures as high as 600 ◦ C. Consequently, the alloys of titanium, specifically Ti–6Al–4V, are found applicable for a spectrum of aerospace application where a combination of weight, strength, and corrosion resistance are required and the high temperature stability of aluminum
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Table 1 Nominal chemical composition of the Ti–6Al–4V alloy (in weight percent). Material
Ti
Al
N
V
C
Fe
H
O
Ti–6Al–4V
90.0
6.0
0.05
4.0
0.1
0.4
0.02
0.20
alloys, high strength steels and the nickel-base super-alloys are insufficient [8]. Since its development and emergence around the late 1940s, the Ti–6Al–4V alloy has grown in strength and preferential selection or choice for use in industries spanning aircraft structural and engine components, biomedical to include surgical implants, and petro-chemical plants due essentially to its excellent combination of mechanical properties, fatigue crack growth resistance and even oxidation resistance. Since its emergence the fatigue behavior of Ti–6Al–4V, particularly fatigue crack growth, has been studied by a number of investigators [9–23]. These studies have essentially shown that (a) microstructure of the alloy to include crystallography and texture [24–27], (b) stress ratio [28–30], and (c) environment [13,31,32] exert a profound influence on fatigue response. For the Ti–6Al–4V alloy categorized as having an alpha plus beta (␣ + ) dual phase, the microstructure is known to exert a dominating influence on fatigue properties. The alpha (␣) phase consists of a hexagonal close-packed (HCP) crystal structure, while the beta () phase is essentially body-centered cubic (BCC). It is generally believed and supported by sufficient experimental data that refining the microstructure helps in an improvement in the fatigue strength of the smooth specimens [9–13]. However, the overall resistance to fatigue crack propagation was found to decrease [14,15]. The dual phase [i.e., alpha plus beta] Ti–6Al–4V alloy when properly heat treated has to offer acceptable combination of strength and ductility when properly heated treated, which makes it stronger than the near-alpha phase alloy and near-beta phase alloy counterpart and a preferred candidate for fatiguecritical applications. In this paper, the results of a recent study aimed at understanding the high-cycle fatigue response and final fracture behavior of samples taken from a welded plate of Ti–6Al–4V alloy are presented and discussed. The as-welded microstructure of the alloy in the region immediately adjacent to the weld was characterized for: (i) the nature and morphology of phases present, and (ii) the presence and distribution of other intrinsic microstructural features. The high-cycle fatigue response of the alloy was determined at a load ratio of 0.1 over a range of maximum stress ( max ) and resultant fatigue life (Nf ). The high-cycle fatigue properties and final fracture behavior of the alloy are discussed in light of the conjoint and mutually interactive influences of maximum applied stress, test specimen orientation, intrinsic microstructural features, and stress-deformation-microstructural feature interactions. Of particular interest in this research study was to systematically evaluate and rationalize the role of test specimen orientation (longitudinal versus transverse) in governing damage initiation, damage propagation and overall cyclic resistance of the alloy microstructure quantified in terms of cyclic fatigue life and final fracture behavior.
ufacturer in the fully annealed condition. The nominal chemical composition of the alloy is provided in Table 1. The tensile and fatigue test specimens were precision machined from both the longitudinal and transverse orientations of the two annealed plates that were welded together using the technique of butt welding. The longitudinal specimens were machined with the major stress axis parallel to the rolling direction of the aswelded titanium alloy plate, while the transverse specimens were machined with the major stress axis perpendicular to the rolling direction of the as-welded titanium alloy plate. The required sections for machining the test specimens were taken from a region of the plate immediately adjacent to the weld (Fig. 1). The test specimens were prepared in conformance with standards outlined in the standard ASTM E-8. At the gage section, the test specimens measured 3.125 mm in diameter and 12.5 mm in length. Final surface preparation was achieved by mechanically polishing the gage section of the test specimens with progressively finer grades of silicon carbide embedded emery paper (320-grit, 400-grit and 600-grit), so as to minimize the effects and/or contributions from surface irregularities, such as, scratches, residual machine marks, and surface finish. 3. Experimental procedures 3.1. Initial microstructure characterization An initial characterization of the microstructure of the aswelded Ti–6Al–4V alloy was done using a low magnification optical microscope. Samples were cut from the welded plate of the Ti–6Al–4V alloy and mounted in bakelite. The mounted samples were then wet ground on progressively finer grades of silicon carbide impregnated emery paper using copious amounts of water both as a lubricant and as a coolant. Subsequently, the ground samples were mechanically polished using five-micron diamond solution. Fine polishing to a perfect mirror-like finish of the surface of each titanium specimen was achieved using one-micron diamond solution as the lubricant. The polished samples were then etched using a reagent that is a solution mixture of 5-ml of nitric acid (HNO3 ), 10 ml of hydrofluoric acid (HF) and 85 ml of water (H2 O). The polished and etched surface of the samples of Ti–6Al–4V were observed in an optical microscope and photographed using a bright field illumination technique.
2. Materials and sample preparation The material chosen for this research study was the widely preferred and chosen alloy Ti–6Al–4V. The alloy was provided by Allegheny Technologies ATI Wah Chang (based in Oregon, USA). Owing to the relative ease of production coupled with good hot workability and receptiveness to heat treatment, the Ti–6Al–4V alloy can be easily solution heat treated and annealed to achieve the desired strength properties. The alloy was provided by the man-
Fig. 1. A schematic of the longitudinal and transverse flat test specimen used for Mechanical testing (tensile and cyclic fatigue).
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Fig. 2. Optical micrographs showing the key micro-constituents, their morphology and distribution, in the Ti–6Al–4V alloy at three different magnifications.
3.2. Tensile and cyclic fatigue tests The tensile properties were determined by uniaxial tensile tests performed on a fully automated, closed-loop servo-hydraulic mechanical test machine [INSTRON-8500 Plus] using a 100 KN load cell. The tests were conducted at room temperature (25 ◦ C) and in the laboratory air environment (Relative Humidity of 55 pct.). The test specimens were deformed in uniaxial tension at a constant strain rate of 0.0001/s. An axial 12.5-mm gage length clip-on type extensometer was attached to the test specimen at the gage section using rubber bands. The stress and strain measurements, parallel to the load line, and the resultant mechanical properties, namely, (i) stiffness, (ii) yield strength, (iii) ultimate tensile strength, (iv) failure stress, and (v) ductility (strain-to-failure), was provided
Fig. 3. Optical micrographs showing minimal influence in microstructure of the heat affected zone and the base metal.
as a computer output by the controller unit of the mechanical test machine. The stress-amplitude controlled high-cycle fatigue tests were performed using a sinusoidal waveform at a stress ratio [R = minimum stress ( min )/maximum stress ( max )] of 0.1. The stress-controlled fatigue tests were conducted at a frequency of 5 Hz. At the chosen stress ratio [R = min / max ], the fatigue tests were conducted over a range of stress amplitudes to establish the variation of maximum stress ( max ) with cyclic fatigue life (Nf ). For the purpose of this investigation the fatigue limit was taken to be the stress below which fatigue failure does not occur even after 106 (one-million) cycles.
3.3. Characterization of damage and analysis of failure The fracture surfaces of the cyclically deformed and failed test specimens of the Ti–6Al–4V alloy were thoroughly examined in a scanning electron microscope (SEM) to determine the macroscopic fracture mode and to concurrently characterize the fine scale topography of the cyclic fatigue fracture surface in order to establish both the macroscopic mechanisms and concurrent microscopic mechanisms governing fracture. The distinction between the macroscopic mode of failure and microscopic fracture mechanisms is based entirely on the magnification level at which the observations are made. The macroscopic mode refers to the overall nature of the failure while the microscopic mechanisms relate to the local failure processes, such as: (i) microscopic and macroscopic void for-
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27 25 – –
Tensile ductility ln(Ao /Af ) (%)
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mation, (ii) growth of the microscopic and macroscopic voids and their eventual coalescence, and (iii) nature, intensity and severity of the fine microscopic cracks and macroscopic cracks dispersed throughout the fracture surface. The samples for observation in the scanning electron microscope were obtained from the failed fatigue specimens by sectioning parallel to the fracture surface.
Reduction in area (%)
4. Results and discussion
148 167 140–158 130–180 921 1029 875–995 800–1100
1020 1153 965–1090 900–1200
ksi MPa
134 149 127–144 115–160
ksi
126 137 105 114
GPa msi
Carpenter technology corporation ASM handbook
Longitudinal Transverse – – Ti–6Al–4V
18 20 15 17
Yield strength Elastic modulus Orientation Material
Table 2 Room temperature tensile properties of Ti–6Al–4V [results are mean values based on duplicate tests].
UTS
MPa
7.8 11.5 16 –
Elongation GL = 0.5 (%)
23.8 21.7 46 –
4.1. Initial microstructure characterization The microstructure of the material under consideration is one of the major factors in determining the mechanical properties like yield strength, tensile strength, ductility, fatigue resistance and resultant fracture behavior. The optical microstructure of the welded Ti–6Al–4V alloy plate is shown in Fig. 1 at three different magnifications. A careful observation of the microstructure over a range of low magnifications in the optical microscope revealed a two-phase duplex microstructure consisting of the HCP equiaxed primary alpha (␣) and interspersed grains containing HCP secondary ␣ and BCC  arranged in a lamellar structure. The primary near equiaxed shaped alpha (␣) grains (light in color) were well distributed in the lamellar matrix of secondary ␣ and transformed beta (dark in color) (Fig. 2). The size of both the alpha (␣) grains and the interspersed grains containing secondary ␣ and  range in size from 10 m to 20 m. The presence of trace amounts of aluminum and oxygen in the Ti–6Al–4V alloy contributes to stabilizing and strengthening of the HCP alpha (␣) phase. This is particularly beneficial for enhancing the hardenability, increasing strength, and improving the response kinetics of the alloy to heat treatment [8]. In Fig. 3 is shown the microstructure of the as-welded alloy plates. 4.2. Tensile properties The room temperature tensile properties of the Ti–6Al–4V are summarized in Table 2. The results are the mean values based on duplicate tests. The elastic modulus, yield strength, ultimate tensile strength, elongation to failure and strength at failure (fracture) were provided as an output of the computer-based data acquisition and test control system. The yield strength was determined by identifying the stress at a point on the engineering stress versus engineering strain curve where a straight line drawn parallel to the elastic portion of the curve at 0.2% offset intersects the curve. The ductility is reported as elongation-to-failure over a gage length of 12.5 mm (0.5 in.). This elongation was measured using a clipon extensometer that was attached to the gage section of the test specimen. The elastic modulus of the test specimens machined immediately adjacent to the weld region of the alloy plate was 126 GPa in the longitudinal (L) orientation and 136 GPa in the transverse (T) orientation. The yield strength was 921 MPa (134 ksi) in the longitudinal (L) orientation, which is approximately 12 percent lower than in the transverse (T) orientation [1029 MPa (149 ksi)]. The ultimate tensile strength is 1020 MPa (148 ksi) in the longitudinal (L) orientation, which is 12 percent lower than in the transverse (T) orientation [1153 MPa (167 ksi)]. The ultimate tensile strength of this material is only marginally higher than the yield strength indicating the occurrence of strain hardening or work hardening beyond yield. Ductility quantified in terms of elongation-to-failure was 7.8% in the longitudinal (L) orientation and 11.5% in the transverse (T) orientation. The reduction-in-area (RA) experienced by the test specimen was as high as 24% in the longitudinal (L) orientation and 22% in the transverse orientation. The yield strength and tensile strength values of the alloy conform well with the val-
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Fig. 4. Variation of maximum stress ( max ) with fatigue life (Nf ) for Ti–6Al–4V.
ues obtained and reported by the manufacturer (ATI Wah Chang) and recorded in ASM Metals Handbook [33]. 4.3. Variation of cyclic stress with fatigue life: high-cycle fatigue properties A gradual degradation in the cyclic strength of a material through the accumulation of damage at the ‘local’ level is known as fatigue. The fatigue life (Nf ) of a structure, or component, is defined as the sum of the number of cycles to crack initiation (Ni ) plus the number of cycles to crack propagation (Np ), which eventually culminates in failure by fracture. A comprehensive study of the cyclic fatigue response of the welded Ti–6Al–4V alloy plates is important for its selection and subsequent use in an array of performancecritical applications. Thus, a goal of this investigation was to extend prior research on high-cycle fatigue of non-fabricated material into an area involving a fabricated structural component in an effort to reflect upon the conditions while in actual use or service. A comprehensive understanding of the fatigue behavior and related damage mechanisms governing failure is rendered difficult by the underlying intrinsic microstructural features and their specific role in governing cyclic deformation and fracture processes during cyclic loading. It has been established in earlier studies that the fatigue properties of the commercially available titanium alloys are governed by the conjoint influences of (a) surface defects, (b) crystallographic texture, and (c) intrinsic microstructural effects to include the following: (i) size and morphology of the interspersed grains with the ␣ and  phases arranged in a lamellar structure, (ii) size of the alpha platelets, and (iii) the presence and role of other intrinsic microstructural features [34–40]. Alloy microstructures that favor a low resistance to crack initiation, even at the low values of applied stress, will usually yield the best endurance during high-cycle fatigue testing and even during tensile resistance. The cyclic fatigue test is the most widely used technique to establish the endurance limit of a chosen metal by essentially determining the variation of maximum stress ( max ) or stress amplitude (/2), with fatigue life as quantified by the number of cycles to failure (Nf ). Metals, alloys and other materials based on metal matrices, which are used at stress levels below the endurance limit can be cycled indefinitely [17]. The endurance limit for this non-ferrous metal is taken to be 106 cycles. The high-cycle fatigue test is the most widely used technique to establish the endurance limit of the chosen metal by essentially determining the variation of maximum stress ( max ) with fatigue life (Nf ). For small, highly stressed components, fatigue life is often controlled by cyclic plasticity at the ‘local’ level coupled with the relatively rapid growth of the fine microscopic cracks through the microstructure of the alloy. For larger components that operate at low stress levels, the fatigue life is initially controlled by micro-
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Fig. 5. Variation of maximum elastic strain ( max /E) with fatigue life (Nf ) for Ti–6Al–4V.
scopic crack formation and subsequently by their growth through the alloy microstructure and eventual coalescence to form one or more macroscopic cracks. The kinetics of crack nucleation in this Ti–6Al–4V alloy is governed by the mutually interactive influences of the following: (i) the rate of cooling, (ii) volume fraction and size of the two key micro-constituent phases, i.e., alpha and beta, and (iii) local stress concentration effects or surface residual stress [17–19]. At the load ratio studied (R = 0.1), the variation of maximum stress ( max ) with cyclic fatigue life (Nf ) reveals increasing fatigue life with a decrease in stress amplitude. This trend is shown by most non-ferrous metals and their composite counterparts. At equivalent maximum stress, the fatigue life of the transverse (T) specimen is marginally greater than the longitudinal (L) counterpart (Fig. 4). In an attempt to understand and rationalize the influence of specimen ductility on high-cycle fatigue response the test data is re-plotted as the variation of maximum elastic strain ( max /E) as a function of cycles-to-failure (Nf ). The maximum elastic strain is taken to be the ratio of maximum stress ( max ) to the elastic modulus (E) of Ti–6Al–4V alloy in that specific orientation. This is shown in Fig. 5. At equivalent values of maximum elastic strain particularly at low values of strain and resultant enhanced fatigue life the longitudinal specimen revealed a marginal improvement in fatigue life over the transverse counterpart (Fig. 6). In an attempt to understand the specific role of alloy microstructure in high-cycle fatigue life at the chosen load ratio of 0.1, the test data is plotted to take into consideration the strength of the Ti–6Al–4V alloy plate in the as-welded condition. The maximum stress is normalized by the yield strength of the alloy plate and reveals minimal difference in fatigue life of the two orientations,
Fig. 6. Variation of the ratio of maximum stress/yield stress ( max / y ) to fatigue life (Nf ) for Ti–6Al–4V.
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Fig. 7. Variation of the ratio of maximum stress/ultimate tensile stress ( max / UTS ) to fatigue life (Nf ) for Ti–6Al–4V.
i.e., longitudinal and transverse of the welded alloy plate, within the limits of experimental scatter. Variation of maximum stress normalized by the ultimate tensile strength of the alloy plate in that specific orientation with fatigue life (Nf ) reveals that at near equivalent values of the ratio of max / UTS , the influence of test
Fig. 8. Comparing the variation of maximum stress ( max ) with fatigue life (Nf ) for Ti–6Al–4V alloy test specimens and ASM Handbook for a given stress ratio of (R = 0.1).
specimen orientation on cyclic fatigue life is minimal at all values of the ratio (Fig. 7). Comparing the fatigue test results obtained in this research study with the stress versus fatigue life response of the alloy documented in the published literature is shown in Fig. 8. This figure reveals the high-cycle fatigue resistance of both the longitudinal
Fig. 9. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Longitudinal) deformed in cyclic fatigue at a maximum stress of 829 MPa and resultant fatigue life (Nf ) of 255,334 cycles, showing: (a) overall morphology of failure. (b) The region of crack initiation and radial outward progression of fatigue damage. (c) The transgranular surface: flat and near featureless in the region of early crack growth. (d) Shallow dimples covering the region of overload reminiscent of locally ductile failure mechanism.
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Fig. 10. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Longitudinal) deformed in cyclic fatigue at a maximum stress of 811 MPa and resultant fatigue life (Nf ) of 674,128 cycles, showing: (a) overall morphology of failure. (b) The region between stable and unstable crack growth showing a healthy population of fine microscopic voids and dimples. (c) High magnification of (b) showing a population of macroscopic and fine microscopic voids. (d) Void coalescence to form a fine microscopic crack and shallow dimples indicative of locally ductile failure mechanisms.
and transverse test specimens taken from the as-welded plate to be noticeably better than the high-cycle fatigue resistance of the as-prepared and annealed Ti–6Al–4V alloy. At equivalent values of maximum stress the improvement was of the order of 100–200 percent. This noticeable difference in cyclic fatigue resistance is ascribed to intrinsic microstructural differences between the conventional alloy and the material chosen for this research study and its contribution to governing cyclic fatigue resistance. 4.4. An analysis and understanding of damage initiation, damage propagation and cyclic fracture The fracture surfaces of the high-cycle fatigue test specimens were carefully examined in a JEOL scanning electron microscope (SEM) at the following levels of magnification: (a) Low magnification in an attempt to identify the nature of microscopic crack initiation, the extent of early crack growth through the alloy microstructure and the kinetics of stable crack growth. (b) Gradually increasing levels of magnification for the purpose of studying the nature of damage initiation, nature and depth of stable crack propagation, and other fine scale features on the fracture surface in the region of overload.
For the Ti–6Al–4V material, the fatigue fracture surfaces revealed only a marginal difference in the topographies at the different levels of maximum stress and resultant fatigue life for both the longitudinal and transverse specimens. On a microscopic scale, the nature and morphology of the fracture surfaces did not vary extensively with maximum stress and fatigue life. Only representative fractographs of the fatigue fracture surfaces of samples of Ti–6Al–4V taken from the welded plate, for both the longitudinal and transverse specimens, at different values of maximum stress and corresponding cyclic fatigue life, are shown in Figs. 9–14.
4.4.1. Longitudinal orientation At a maximum stress of 829 MPa, the longitudinal specimen had a fatigue life of 255,334 cycles. Low magnification observations in the scanning electron microscope revealed the existence of well defined and distinct region of crack initiation and a short region representative of stable crack propagation (Fig. 9a). Early microscopic crack growth was observed to be extending radially outward from the point marking the onset of crack initiation (Fig. 9b). A careful examination of the region of stable crack propagation at the higher allowable magnifications of the SEM revealed the region to be flat and featureless, containing a population of fine microscopic cracks that were randomly distributed through the region (Fig. 9c).
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Fig. 11. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Longitudinal) deformed in cyclic fatigue at a maximum stress of 811 MPa and resultant fatigue life (Nf ) of 674,128 cycles, showing: (a) The transgranular surface in the region of early crack growth. (b) High magnification of (a) showing pockets of fine shallow striations indicative of localized micro-plastic deformation. (c) Well defined striations in the region of stable crack growth indicative of micro-plastic deformation.
For almost all the samples cyclically deformed at varying levels of maximum stress, the overload region revealed the presence of a population of shallow dimples indicative of locally ductile failure mechanisms (Fig. 9d). For the test specimen that was cyclically deformed at a maximum stress of 811MPa the resultant fatigue life was 674,128 cycles. The fracture surfaces of this test specimen revealed distinctly the regions of crack initiation and overload (Fig. 10a). In the terminal stages of the region of stable crack growth through the alloy microstructure, a transition region between stable and unstable crack growth was distinctly evident (Fig. 10b). High magnification observations of the transition region revealed a population of macroscopic voids and fine microscopic voids (Fig. 10c). Continued cyclic loading led to a progressive growth of the fine microscopic voids and their eventual coalescence resulting in the formation of one or more fine microscopic cracks. Half of these microscopic voids are the shallow dimples visible on the fracture surface (Fig. 10d). The transgranular region (Fig. 11a), when observed at high magnifications, revealed (a) pockets of fine shallow striations in the region of early microscopic crack growth, and (b) well defined striations (Fig. 11b). This feature is indicative of the localized micro-plastic deformation mechanisms in the region of stable crack growth. Thus, for the longitudinal orientation, the microscopic fracture characteristics observed for the test sample that was cyclically
deformed at high maximum stress, resultant low fatigue life, and for the test sample that was cyclically deformed at a low maximum stress and resultant enhanced fatigue life, are essentially similar. 4.4.2. Transverse orientation Following the same trend as the specimens oriented in the longitudinal direction, the test specimens prepared from the transverse orientation of the as welded Ti–6Al–4V alloy plate exhibited almost identical fatigue lives at the higher level of maximum stress. At a maximum stress of 962 MPa and resultant fatigue life of 40,664 cycles, the fracture surface features are shown in Fig. 12. The overall morphology comprised of (a) distinct regions of crack initiation and growth, (b) a short region of stable crack growth, and (c) overload (Fig. 12a). The region of early crack growth revealed radial propagation of the fatigue damage from the onset of initiation (Fig. 12b). High magnification observations in the region of stable crack growth revealed fine and shallow striations indicative of localized micro-plastic deformation (Fig. 12c). The region of overload revealed a population of both fine microscopic voids and macroscopic voids. For a specimen subjected to a maximum stress of 900 MPa and resultant enhanced fatigue life of 425,984 cycles, the overall fracture surface morphology revealed distinctive regions of crack initiation and radial propagation of the cracks across the transgranular surface region (Fig. 13a). High magnifica-
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Fig. 12. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Transverse) deformed in cyclic fatigue at a maximum stress of 962 MPa and resultant fatigue life (Nf ) of 40,664 cycles, showing: (a) overall morphology of failure. (b) The region of early crack growth: flat and near featureless. (c) Pockets of shallow striations in the region of stable crack growth. (d) A population of fine microscopic and macroscopic surface voids in the region of overload.
tion observations in the region of early crack propagation (Fig. 13b) revealed a random distribution of fine microscopic cracks (Fig. 13c). The characteristic features on the overload surface were fine microscopic voids and a population of shallow dimples, and these were observed in test specimens that were cyclically deformed at both high maximum stress and low maximum stress (Figs. 12d and 13d). Repeated pockets of very fine and shallow striations were found at random locations in the region of stable crack growth indicative of the highly localized ductile failure mechanisms (Fig. 14). 4.5. The mechanisms governing the initiation of damage, propagation of damage and fracture at microscopic level It is well known that under cyclic stress control, the resultant fatigue life is dependent on (a) the nature of crack initiation, and (b) the nature and severity of crack propagation through the alloy microstructure. Much of the fatigue life is spent or utilized for the initiation of fine microscopic cracks. The intrinsic microstructural features that were formed during initial heat treatment and during subsequent fabrication of the welded alloy plate do exert an influence on cyclic fatigue response and resultant fatigue life of the candidate alloy. The microstructure of the Ti–6Al–4V alloy material chosen for this study is taken from a region, which can be characterized as having a duplex microstructure consisting of near equiaxed alpha (␣) and transformed beta () phases. In order to
account for requirements of continuity, slip in the HCP primary ␣ phase accounts for majority of the plastic deformation in this alloy. The lamellar structure merely acts to inhibit the slip in the primary ␣ phase and this occurs at the ␣- phase boundaries. An earlier study on ␣- alloy titanium alloy has shown that slip initiates in the primary ␣ phase, which is mechanically softer than the  lamellar phase. Further, the mechanical properties of this and other two phase alloys of titanium are very sensitive to the geometrical arrangement of the phases and the crystallographic texture. Other companion and subsequent studies by researchers have shown that fatigue cracking in the Ti–6Al–4V alloy initiates both at and along the grain boundaries. This was attributed to the presence of a large area fraction of weakened high angle grain boundaries parallel to the longitudinal direction. The preference for crack initiation and early crack propagation along the grain boundaries results in a progressive loss of the through thickness constraint causing an essentially plane strain fracture process to be divided into several plane stress fractures. During continued cyclic deformation of the alloy specimens, a gradual degradation of the intrinsic microstructural features results in the formation of “localized” failure sites coupled with the accumulation of strain at the microscopic level. When the strain accumulation and resultant stress concentration at the localized level exceeds a critical value, microscopic crack initiation is favored to occur, which is exacerbated by the continued application of stress culminating
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Fig. 13. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Transverse) deformed in cyclic fatigue at a maximum stress of 900 MPa and resultant fatigue life (Nf ) of 425,984 cycles, showing: (a) overall morphology of failure showing region of initiation and radial progression of fatigue damage. (b) High magnification of region of early crack growth. (c) High magnification observation of the region of early crack growth showing a random distribution of fine microscopic cracks. (d) A population of voids and dimples in the region of overload.
in the progressive accumulation of micro-plastic deformation at the ‘local’ level. Continuous cyclic deformation ably aided by the continuous propagation of the fine microscopic cracks resulting from the growth and eventual coalescence of the microscopic voids leads to the formation of macroscopic cracks. A gradual growth or propagation of both the fine microscopic and macroscopic cracks through the alloy microstructure results in a progressive drop in load-carrying capability of the test specimen culminating in failure. Based on a careful and comprehensive observation of the fracture surfaces of the test specimens taken from both the longitudinal and transverse orientations of the as-welded titanium alloy plates and cyclically deformed at specific values of maximum stress, the fracture plane of the cyclically deformed and failed test specimens was essentially normal to the stress or loading axis, suggesting the importance of tensile stress in promoting fracture. Few of the microscopic voids that were created by the presence of intrinsic microstructural features did not grow extensively in the direction of applied stress, which is generally the case for ductile failure of metallic materials. The lack of extensive growth of the fine microscopic voids suggests that the overall strain that induces fracture at both the microscopic and macroscopic level is controlled by (i) the strain required for void nucleation, (ii) crack propagation strain, and (iii) the eventual strain required for their linkage.
Fig. 14. Scanning electron micrographs of the fatigue fracture surface of the Ti–6Al–4V specimen (Orientation: Transverse) deformed in cyclic fatigue at a maximum stress of 823 MPa and resultant fatigue life (Nf ) of 581,951 cycles, showing pockets of shallow striations reminiscent of local micro-plastic deformation.
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5. Conclusions A study of the high-cycle fatigue and final fracture behavior of test specimens machined from a Ti–6Al–4V as welded plate for the two different orientations, i.e. longitudinal and transverse provides the following key observations: 1. The Ti–6Al–4V alloy exhibited a duplex microstructure consisting of the near equiaxed alpha (␣) and the interspersed grains of ␣ and  phases. The primary near equiaxed shaped alpha (␣) grains was well distributed in a lamellar matrix with transformed beta. 2. The yield strength and ultimate tensile strength of Ti–6Al–4V is nearly equal both in the longitudinal orientation and in the transverse orientation indicating a more or less isotropic nature of the material. In both the longitudinal and transverse orientations, the tensile strength of the as-welded alloy plate is higher than the yield strength indicating the occurrence of work hardening beyond yield. 3. There is only a marginal influence of microstructure on highcycle fatigue life of the as-welded alloy plate in the two orientations. The variation of maximum stress with cyclic fatigue life (Nf ) revealed a trend shown by most non-ferrous metallic materials. At equivalent values of maximum stress, at a load ratio of 0.1, the fatigue life of the longitudinal and transverse specimens show minimal difference at all levels of stress, with the transverse specimen exhibiting marginally improved fatigue life when compared to the longitudinal counterpart. 4. At the chosen load ratio of 0.1 the fatigue fracture surfaces revealed minimal to no difference in the nature, distribution and volume fraction of the intrinsic features on the fracture surface as a function of maximum stress and resultant fatigue life. The region of crack initiation and early microscopic crack growth and a short, yet distinct, region of stable crack growth were essentially flat and transgranular. A population of fine microscopic and macroscopic cracks was evident in the region of stable crack growth approaching overload. Also evident in this region were pockets of well defined striations indicative of the occurrence of micro-plastic deformation at the local level. The region of overload was covered with a population of shallow dimples, fine microscopic voids and isolated macroscopic void and cracks, features reminiscent of the locally governing failure mechanisms. During repeated cyclic loading, the fine microscopic cracks coalesce to form macroscopic crack. References [1] M. Peters, R. Braun, 4th European Conference on Advanced Materials and Processes – EUROMAT 95, Associazione Italiana di Metallurgia, Milano, 1995, p. 55.
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