Influence of microstructure on strain-controlled fatigue and fracture behavior of ultra high strength alloy steel AerMet 100

Materials Science & Engineering A 601 (2014) 29–39

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Influence of microstructure on strain-controlled fatigue and fracture behavior of ultra high strength alloy steel AerMet 100 K. Manigandan a, T.S. Srivatsan a,n,1, Deepthi Tammana b, Behrang Poorganji b, Vijay K. Vasudevan b a b

Department of Mechanical Engineering, The University of Akron, Akron, OH 44325-3903, USA Department of Mechanical & Materials Engineering, University of Cincinnati, 2901 Woodside Dr, Room 501B ERC, Cincinnati, OH 45221-0072, USA

art ic l e i nf o

a b s t r a c t

Article history: Received 19 November 2013 Received in revised form 29 January 2014 Accepted 30 January 2014 Available online 6 February 2014

In this paper, the results of a study aimed at understanding the specific role of microstructure on cyclic stress response, cyclic strain resistance, and cyclic stress versus strain response, deformation and fracture behavior of high strength alloy steel AerMets 100 is presented and discussed. The cyclic strain amplitude-controlled fatigue properties of this ultra-high strength alloy steel revealed a linear trend for the variation of log elastic strain amplitude with log reversals-to-failure, and log plastic strain amplitude with log reversals-to-failure. Cyclic stress response revealed a combination of initial hardening for the first few cycles followed by stability for large portion of fatigue life before culminating in rapid softening to failure at the lower cyclic strain amplitudes and intermediate cyclic strain amplitudes and resultant enhanced cyclic fatigue life. Fracture characteristics of test specimens of this high strength alloy steel were different at both the macroscopic and fine microscopic levels over the entire range of cyclic strain amplitudes examined. Both macroscopic and fine microscopic observations revealed fracture to be essentially ductile with features reminiscent of locally occurring ductile mechanisms. The intrinsic microscopic mechanisms governing stress response, deformation characteristics, fatigue life and final fracture behavior are presented and discussed in light of the competing and mutually interactive influences of intrinsic microstructural effects, deformation characteristics of the microstructural constituents, cyclic strain amplitude and concomitant response stress. & 2014 Elsevier B.V. All rights reserved.

Keywords: Alloy steel Microstructure Cyclic strain amplitude Stress response Fatigue life Deformation and fracture

1. Introduction Sustained research and development efforts have in recent years culminated in the emergence of the family of high strength steels having a dual-phase and even triple-phase microstructure [1–8]. This has been made possible through novel innovations in the domains spanning chemical composition, primary processing, secondary processing and even manufacturing techniques. For example, a careful control of heating and subsequent mechanical deformation, by either rolling or thermo-mechanical processing (TMP), of the alloy steels can result in the formation of either a dual-phase microstructure or a triple-phase microstructure that has a fine grain size and offers high strength. A suitable combination of the alloying elements, such

n

Corresponding author. E-mail addresses: [email protected] (T.S. Srivatsan), [email protected] (V.K. Vasudevan). 1 Member of Editorial Board of MSE A. http://dx.doi.org/10.1016/j.msea.2014.01.094 0921-5093 & 2014 Elsevier B.V. All rights reserved.

as, nickel, copper, cobalt and molybdenum does contribute to increasing strength of the alloy steel by the following two approaches: (i) Directly through a synergism of solid solution strengthening, precipitation hardening, and overall refinement of the intrinsic microstructural features. (ii) Indirectly through enhanced hardenability and associated modification of the microstructure [9].

A noticeable advantage in selecting and using a high strength alloy steel for a chosen application is that they offer the possibility of reducing overall weight of the component or structure thereby enabling in achieving [10,11]: (i) a reduction in overall cost of fabrication, (ii) a direct reduction in cost of transportation of the component or structure, (iii) facilitating ease in handling of the material or component, and (iv) overall better surface finish over other high strength alloy steels having near similar chemical composition. Furthermore, light weight and thin components and structures are both essential and desired for reducing the

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K. Manigandan et al. / Materials Science & Engineering A 601 (2014) 29–39

cost of operation of the engineering component besides enabling in the possibility of creative designs for the structure and/or component. Also, an observable reduction in size of the end product, or component, translates into less consumption of the “candidate” steel during manufacturing. Increased interest in the family of high strength alloy steels has resulted from the many advantages they have to offer particularly when chosen as the candidate for use in performance-critical applications. This has provided the much desired interest and incentive to enable continuing studies on steels with the primary purpose of understanding microstructural influences on their mechanical properties to include fatigue resistance, thereby evaluating their endurance while in service. The steel AerMet 100 was developed and emerged as a viable replacement to 300 M steel the material that was preferred and used for landing gear of the carrier-based aircraft. This steel was found to be preferentially better than 300 M steel due on account of its greater fracture toughness and overall better resistance to stress corrosion cracking and hydrogen embrittlement [12]. In the early 1990s E.U. Lee at the Naval Air Development Center investigated the fatigue crack growth behavior of this steel under conditions of constant amplitude loading in both inert and corrosive environments [13]. In both environments, the fatigue crack growth rates above 10
6 mm/cycle were found to be essentially unaffected by the load ratio. However, at near threshold levels it was observed that increasing the stress ratio resulted in a reduction in the threshold stress intensity range for fatigue crack growth. Also, at low growth rates, i.e., below 10
6 mm/cycle the crack growth rate was observably greater and the threshold stress intensity factor (ΔKTH) was smaller in dry nitrogen gas than in 3.5 pct. NaCl solution for all of the stress ratios studied, i.e., 0.1, 0.5 and 0.8. The properties of a high strength alloy steel are governed by the mutually interactive influences of chemical composition, processing history, intrinsic microstructural features, temperature, nature of loading and even the test environment to which the material is exposed to while in service [14–20]. Conditions that are favorable to minimizing constraints at the tip of a crack-like defect are fully conducive for enhancing the ability of the steel to plastically deform. This not only facilitates in achieving an overall improvement in fracture toughness but also provides improved tolerance to damage due to fatigue. To totally prevent premature failure of the structure or component and to concurrently improve its fatigue life it is essential to determine the fatigue behavior of the steel of interest by mechanical testing. By studying both the initial microstructure and deformed microstructure of the chosen material, i.e., high strength alloy steel, it is possible to establish both the macroscopic and fine microscopic mechanisms influencing the fatigue behavior of this steel. In this specific study, the fatigue behavior of AerMet 100 was determined under conditions of strain control at room temperature (T ¼27 1C). Variation of strain with fatigue life was determined using the Basquin equation and the Coffin–Manson equation [21,22]. Variation of the monotonic (static) stress with strain and cyclic stress with strain was determined using the Ramberg–Osgood equation [23,24]. The underlying mechanisms governing cyclic strain amplitude-controlled deformation and eventual fracture behavior of the chosen alloy steel are discussed in light of the mutually interactive influences of intrinsic microstructural effects, nature and magnitude of loading, deformation characteristics of the

microstructural constituents and fatigue life. On this alloy steel little to no work has been done and documented in the published literature on aspects related to strain-controlled low cycle fatigue, deformation characteristics and fracture behavior. Specific details pertaining to the stress controlled high cycle fatigue behavior of the chosen steel can be found elsewhere [25]. Hence, the motivation for this exhaustive research study, the results and findings of which are described in this technical paper.

2. Material The test material chosen for this experimental study is AerMet 100 that belongs to the family of developed and emerged high strength alloy steels fort use in aircraft landing gears. The material was provided by Carpenter Technology Corporation (Reading, PA, USA). The material [UNSK # 9258] was produced through use of the methods of vacuum induction melting (VIM) and vacuum arc remelting (VAR) and then immediately cast into ingots. The ingot was then hot worked to get blocks that measured 19 mm square. The blocks were then hot rolled to get round bars which were 19 mm in diameter. The hot rolled bar (measuring 19 mm in diameter) was subsequently annealed at 1275 F (677 1C) for 6 h and then turned and ground to size. The ground bar was tested for defects using immersion sonic inspection. The hardness of the bar in the annealed condition was 40HRC. This alloy steel is particularly noted for offering an attractive combination of high hardness and high strength coupled with acceptable ductility and adequate toughness. This alloy steel has been chosen for use in components that require both strength and fracture toughness coupled with exceptional resistance to stress corrosion cracking (SCC) and fatigue loading [12]. The nominal composition of the starting material (in weight percent) is given in Table 1.

3. Experimental procedures 3.1. Preparation of test specimen Cylindrical test specimens, conforming to specifications outlined in ASTM E8 [26] were precision machined from the as-provided round bars (0.75 in. diameter) of the chosen steel. The threaded test specimens measured 114.3 mm in length and 12.70 mm in diameter at the thread section. The gage section of the machined test specimen measured 25.4 mm in length and 6.35 mm in diameter. The ratio of overall length to diameter of the fatigue test specimen was chosen to ensure that it would not buckle even under fullyreversed total strain amplitude-controlled cyclic deformation. A schematic of the test specimen is shown in Fig. 1. To minimize the effects and/or contributions arising from surface irregularities and surface finish, final preparation of the test specimen surface was achieved by mechanically polishing the gage section of all specimens to remove any and all circumferential scratches and surface machining marks. 3.2. Microstructural characterization Samples of the steel taken from both the longitudinal (L) and transverse (T) orientations were prepared very much in conformance

Table 1 Nominal Chemical composition of AerMet 100 (in weight percent). Material

C

Mn

Si

P

S

Cr

Ni

Mo

Co

Al

Ti

AerMet 100

0.238

o 0.01

0.03

0.002

0.0007

2.99

11.2

1.18

13.4

0.003

0.011

K. Manigandan et al. / Materials Science & Engineering A 601 (2014) 29–39

25.4

25.4

6.35

12.7 7.62 25.4

114.3 Fig. 1. A schematic of the cylindrical test specimen used for mechanical testing.

with standard procedures used for the metallographic preparation of metal samples. This involved coarse polish using progressively finer grades of silicon carbide (SiC) impregnated emery paper [i.e., 320grit, 400-grit and 600-grit] followed by fine polishing using 5 μm and 1 μm alumina-based polishing compound suspended in distilled water as the lubricant. The as-polished samples were than etched using nital reagent, i.e., a solution mixture of nitric acid in methanol. Etching helps in revealing the grain boundaries, morphology of the grains, and other intrinsic features in the microstructure. The polished and etched samples were examined in an optical microscope, at low magnifications, and photographed using standard bright field illumination technique. Transmission electron microscopy (TEM) observations were used to investigate the following: (i) Grain size, grain distribution, presence and distribution of second-phase carbide particles and other fine intrinsic features in the microstructure of the as-provided material. (ii) Nature and distribution of the deformation structures in the cyclically deformed and failed test specimens. Samples of 0.5–0.6 mm in thickness were sliced from both the undeformed section, i.e. grip section, and deformed section (i.e. gage length) of the alloy steel test sample. From these slices was punched out discs having an effective diameter of 3 mm and a thickness of around 100 μm. The circular discs, referred to henceforth as samples, were dimpled using a dimpler machine [Model: FISCHIONE] using diamond solutions gradually down to 1 μm size. Using the technique of dimpling, the thickness of the sample was effectively reduced to around 20–30 μm. Subsequently, ion milling was used for the final stage of sample preparation for observation in a transmission electron microscope (TEM). Argon gas ion milling was used with gun angle set at 141 and in a chamber that was cooled by liquid nitrogen. The resultant polished sample, referred to as “thin-foil”, was observed in a transmission electron microscope (Model: Phillips) operating at 200 kV. 3.3. Mechanical testing Uniaxial tensile tests were performed up until failure on a fullyautomated, closed-loop servohydraulic mechanical test machine [INSTRON Model 8500 plus] equipped with a 10, 000 kgf (98 KN) load cell. The test specimens were deformed at a constant strain rate of 0.0001 s
1. An axial 12.5 mm gage length extensometer was attached to the test specimen at the gage section, using rubber bands, to provide a precise measurement of strain during uniaxial loading and resultant stretching of the test specimen. The stress and strain measurements, parallel to the load line, were recorded on a PC-based data acquisition system [DAS]. The cyclic fatigue tests were also performed on a fully-automated closed-loop servo-hydraulic structural test machine. The tests were conducted in the axial total strain amplitude control mode under fully reversed, push–pull, tension–compression loading. For each individual test the test machine was programmed to maintain a

31

constant nominal strain rate of 0.001 s
1. The test common signal (strain function) was a triangular waveform, and the mean strain was zero. All total strain amplitude-controlled (ΔεT/2) cyclic fatigue tests were initiated in tension. An axial 12.7 mm clip-on extensometer was attached to the test specimen at the gage section, using rubber bands, for the purpose of monitoring total strain–amplitude during fullyreversed strain amplitude-controlled fatigue tests such that the magnitude of negative strain equals the magnitude of positive strain (Rε ¼mminimum/εmaximum ¼
1). The controlled variable is total strain amplitude (ΔεT/2) and the cyclic strain-controlled tests were performed at six different total strain amplitudes (ΔεT/2). The extensometer was calibrated prior to the initiation of each fatigue test. The tests were conducted in laboratory air environment (temperature of 27 1C and relative humidity of 55%). The stress and strain data for each fatigue test were recorded on a PC-based data acquisition system. The number of cycles-to-failure, i.e., separation of the test specimen into two parts, is taken as fatigue life (Nf). During total strain amplitude-controlled cyclic deformation a measurable physical quantity is plastic strain. Cyclic plastic strain does tend to produce a number of damaging processes, which affect the microstructure, cyclic stress and strain response and resultant low-cycle fatigue life. The damage resulting from cyclic straining helps relate the fatigue life (Nf) to strain amplitude (Δε/2) in a fullyreversed [Rε ¼
1] strain amplitude-controlled fatigue test. 3.4. Failure-damage analysis Fracture surfaces of the cyclically deformed and failed fatigue test specimens of this alloy steel were comprehensively examined in a scanning electron microscope (SEM) to (a) Determine the macroscopic fracture mode, and concurrently (b) Characterize: (i) the fine scale topography, (ii) the nature of crack initiation, (iii) the extent and depth of early crack propagation, (iv) the extent and depth of stable crack propagation, and (v) other intrinsic features on the fracture surface. This was essential for the purpose of establishing the microscopic mechanisms contributing to failure by fracture. The distinction between the macroscopic mode and microscopic fracture mechanisms is based entirely on the magnification level at which the observations were made. The macroscopic mode refers to the overall nature of failure while the fine microscopic mechanisms relate to the failure processes occurring at the “local” level, to include the following: (i) microscopic void formation, (ii) microscopic void growth and their eventual coalescence by way of impingement, and (iii) nature, intensity and severity of the fine microscopic and macroscopic cracks dispersed through the fracture surface. The samples for observation in the scanning electron microscope (SEM) were obtained from the failed low cycle fatigue specimens by sectioning parallel to the fracture surface.

4. Results and discussion 4.1. Microstructure The optical microstructure of the as-provided alloy steel is shown in Fig. 2. The observed microstructure is quite typical of high strength alloy steel in that it clearly reveals a combination of carbon-rich, identified as the darker regions in the microstructure and carbon-depleted regions, identified as the light-color regions in the microstructure. When compared to the widely chosen and used alloy steels, a higher carbon [C¼0.238 pct.] coupled with alloy content [Cr¼ 2.99 pct., Ni¼11.2 pct., Mo¼1.18 pct., Co¼ 13.4 pct.] in this steel resulted in a greater volume fraction of the martensite

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Ferrite

Martensite

10μm

5μm

Fig. 2. Optical micrographs showing the key micro-constituents present in alloy steel AerMets 100, at two different magnifications.

200n

200n

Fig. 3. Bright field transmission electron micrographs showing the elongated nature of grains in the alloy steel in the undeformed condition.

Table 2 A compilation of the room temperature tensile properties of AerMet 100. Specimen

AerMet 100

Elastic modulus

Yield strength

UTS

ksi

GPa

ksi

MPa

ksi

MPa

275,340

190

255

1760

289

1993

micro-constituent in the carbon-rich regions. The martensite was easily noticeable and present in large numbers in the form of “finelath” in this steel that was produced through a synergism of vacuum induction melting (VIM) and vacuum arc re-melting (VAR) prior to casting to get an ingot. The cast ingot was annealed and provided for this research study. The presence of the martensite micro-constituent in the carbon-rich regions is governed by a synergism of both composition of the chosen steel (AerMet 100) and primary processing technique used to get the starting product, and does exert an influence on the following: (i) tensile properties, (ii) fatigue properties to include stress response characteristics and fatigue life, and (iii) final fracture behavior. Transmission electron microscope observations of the thin foil taken from the undeformed alloy steel when viewed at low magnifications revealed the grain to be small in diameter and elongated. Essentially, the grains were found to be of varying size and orientation (Fig. 3). No attempt was made to measure the

Elongation GL ¼ 1″ (%)

Reduction in area (%)

Tensile ductility ln(Ao/Ar) (%)

28

66

105

grain size and thereby determine the average grain size of the “undeformed” material.

4.2. Tensile properties The tensile properties, at ambient temperature (27 1C) are summarized in Table 2. Results reported are the mean values based on duplicate tests. The elastic modulus of the steel is 190 GPa. The yield strength of this steel is 1760 MPa, while its ultimate tensile strength is 1993 MPa. The tensile strength (sUTS) is only marginally higher than the yield strength (sYS) indicating low value of strain hardening beyond yield. The ductility, quantified by elongation over 0.5 in. (12.7 mm) gage length, was as high as 28%. The reduction in test specimen cross-sectional area, another measure of ductility was as high as 66%. The tensile ductility, defined as ln Ao/Af, was as high as 105%. The variation of engineering stress with engineering strain

K. Manigandan et al. / Materials Science & Engineering A 601 (2014) 29–39

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Table 3 Summary of properties obtained from strain-controlled fatigue tests. Δ A T/2 ΔA e/2 (%) ΔA p/2 (%) Δs/2 (MPa) (%) N ¼ 1 N ¼ Nf/2 N¼ 1 N ¼Nf/2 N ¼ 1 N ¼Nf/2

Specimen

AerMet 100 1.2 1.05 0.9 0.7 0.6 0.5

10 Cyclic Strain Amplitude (%)

AerMet® 100 T=27°C

1 b=-0.06

Cyclic Stress Amplitude (MPa)

2000 Fig. 4. The engineering stress versus engineering strain curve obtained for alloy steel AerMets 100 at room temperature.

0.95 0.78 0.86 0.64 0.58 0.48

0.84 0.74 0.81 0.62 0.58 0.49

0.25 0.27 0.04 0.06 0.03 0.03

0.36 0.31 0.09 0.08 0.02 0.02

1914 1879 1710 1443 1322 1034

Nf cycles

1729 1735 1638 1468 1339 1040

65 113 692 1266 3634 10,000

AerMet® 100 T=27°C

Δ

1800 1600

1.05

1400 1200

0.9 0.8

0.7

1000 1.2

0.55

800

f’=1.97%

100

0.1

101

102

103

104

105

Fatigue Life (Nf) Plastic Elastic Total

0.01 101

c=--0.15

102

103

Fig. 6. Cyclic stress response curves for AerMets 100 over the entire range of strain amplitudes, when deformed at room temperature (27 1C).

104

Reversals to Failure (2Nf) Fig. 5. Variation of cyclic strain amplitude (Δε/2) with reversals-to-failure (2Nf) for alloy steel AerMets 100 at room temperature (27 1C).

is shown in Fig. 4. This steel revealed evidence of large plastic strain prior to failure by fracture.

At the six total strain amplitudes the alloy steel was cyclically deformed it is observed that at both N ¼1 and N ¼ Nf/2 (half-life) an observable portion of the total strain amplitude was elastic strain amplitude [Refer to Table 3]. Similarly, on a log–log plot the variation of plastic strain amplitude [Δεp/2] with reversals-tofatigue life (2Nf) is also linear and the Coffin–Manson relationship is satisfied. ½Δεp =2 ¼ ε0f ½ð2Nf Þc 

4.3. Low-cycle fatigue (LCF) resistance. The effect of fully reversed strain cycling on low-cycle fatigue response of this high strength steel is as shown in Fig. 5. The cyclic strain amplitude [Δε/2] versus reversals-to-fatigue life (2Nf) curve can be viewed as an indication of the resistance of the alloy steel microstructure to both crack initiation and eventual failure by fracture. Throughout LCF testing at the different cyclic strain amplitudes no evidence of specimen buckling was observed. The elastic strain amplitude and plastic strain amplitude were determined from the horizontal span of the hysteresis loop for each cycle. The cyclic stress amplitude was determined from the vertical span of the hysteresis loop for each cycle recorded by the control unit of the servohydraulic test machine [Model: INSTRON 8500 Plus]. On a log–log scale the variation of elastic strain amplitude [Δεe/2] with reversals-to-fatigue life (2Nf) is linear and the Basquin relationship is satisfied ½Δεe =2 ¼ ðs0f =EÞ ð2N f Þb 

ð1Þ

In this equation [Δεe/2] is the elastic strain amplitude, E is the elastic modulus of the steel, 2Nf is the number of reversals to fatigue failure and s0f is the fatigue strength coefficient while b is the fatigue exponent.

ð2Þ

In this equation Δεp/2 is the plastic strain amplitude, 2Nf is the number of reversals to failure, ε0f is the fatigue ductility coefficient, and c is the fatigue ductility exponent. These two equations [i.e., Eqs. (1) and (2)] can now be used to determine the cyclic parameters [b¼ 0.06 and c ¼0.15] in the low-cycle fatigue regime [Nf o104 cycles] [21–24]. 4.4. Cyclic stress response characteristics An important feature of the strain-controlled fatigue process is the variation of stress response with fully-reversed strain cycling. The stress response curves were established by monitoring the cyclic stress range during fully-reversed total strain amplitude-controlled fatigue. These curves provide useful information pertaining to overall cyclic stability of alloy steel AerMets 100. Overall stability of the intrinsic microstructural features, during fully-reversed strain cycling, coupled with an intrinsic ability of the microstructure of alloy steel AerMet 100 to distribute plastic strain over the entire microstructural volume and among the different microstructural constituents are important factors controlling: (i) stress versus strain response, (ii) cyclic strain resistance, and (iii) fatigue life. The stress response curves for this high strength steel, over a range of total strain amplitudes, are shown in Fig. 6. Since plastic strain amplitude varied throughout testing, reaching a maximum

K. Manigandan et al. / Materials Science & Engineering A 601 (2014) 29–39

value at the minimum stress and as minimum value at the maximum stress, the value at specimen half-life (Nf/2) is taken as the reference. At low cyclic strain amplitudes and intermediate cyclic strain amplitudes this high strength alloy steel showed evidence of initial hardening during the first few cycles of fullyreversed loading followed by stability for most part of fatigue life and culminating in rapid softening in the few cycles prior to catastrophic failure. However, at a higher cyclic strain amplitude, i. e., ΔεT/2 ¼or 41.05%, this steel showed evidence of gradual softening to failure from the onset of fully-reversed cyclic deformation. The rapid and observable decrease in stress carrying capability prior to failure can be ascribed to the occurrence of the following mutually interactive events, based on examination of the fatigue fracture surfaces in the SEM as shown and described in the following section on cyclic fatigue fracture: (a) A gradual degradation of the microstructure through the formation and presence of fine microscopic cracks. (b) A progressive growth of the fine microscopic cracks through the microstructure during continued cyclic straining. (c) Eventual coalescence or linkage of the fine microscopic cracks to form one or more macroscopic cracks. Prior to failure the softening effect is exacerbated by the concurrent growth of both the fine microscopic and macroscopic cracks through the microstructure of alloy steel AerMet 100. Variation of normalized stress (sN/s1), i.e., defined as the ratio of stress at any cycle (sN) with respect to stress in the first cycle (s1) with cycles, as shown in Fig. 7, clearly reveals the softening effect, i.e., an observable decrease in stress response with fullyreversed strain cycling, to be far more pronounced and noticeable at the higher cyclic strain amplitudes and resultant short fatigue life than at the intermediate cyclic strain amplitudes and lower cyclic strain amplitudes and resultant enhanced fatigue life. During fully-reversed strain cycling of high strength alloy steel AerMet 100 the loss in strength arising from the simultaneous and interactive influences of (a) intrinsic microstructural changes, and (b) the occurrence of both microscopic and macroscopic cracking, by far offsets the increased strength that results as a result of (i) dislocation–dislocation interactions, and (ii) dislocation-secondphase particle interactions. Variation of cyclic stress amplitude (Δs/2) with percentage life (N/Nf) during fully-reversed strain cycling is shown in Fig. 8. This figure reveals the degree of softening or hardening experienced by the chosen alloy steel as a fraction of total fatigue life. The cyclic response of this high strength alloy steel AerMets 100 is noticeably different at the high cyclic strain amplitudes, intermediate cyclic strain amplitudes and low cyclic strain amplitudes. Rapid initial softening to failure was evident at 1.04

Δ

AerMet® 100 T=27°C

T/2(%)=

Normalized Stress σN/σ1

1.02 1.00 0.98

1.05

0.8

0.96 0.9

0.55

0.94

0.7 1.2

2200 AerMet® 100 T=27°C

Δ

Cyclic Stress Amplitude (MPa)

34

2000 1800 1600

1.05 0.9

1400 0.7

1200

0.8 1.2

1000

0.55

800 0

20

40

60

80

100

120

Percentage Life (N/Nf) Fig. 8. Variation of cyclic stress amplitude (MPa) with percentage life (N/Nf) over the entire range of cyclic strain amplitudes investigated.

the higher cyclic strain amplitudes. At the intermediate cyclic strain amplitudes (i.e., Δεp/2¼ 1.05%, 0.90% and 0.80%) rapid initial softening during the first twenty percent of life was followed by gradual, yet noticeable, softening to failure. At the lower plastic strain amplitudes [i.e. Δεp/2¼0.7 and 0.55] the degree of softening was minimal prior to culminating in failure. At all values of cyclic strain amplitude the compressive stress is lower than the tensile stress from the onset of fully-reversed cyclic deformation indicating the occurrence of Bauschinger effect. This behavior is very much in conformance with the trend expected for high strength metals, such as high strength steels and age hardenable aluminum alloys, during fully-reversed cyclic straining. For most of fatigue life, stress response comprised of a gradual decrease of both the tensile stress (sT) and compressive stress (sC) components at the high cyclic strain amplitude [ΔεT/2] and resultant higher plastic strain amplitude [Δεp/2]. However, at the intermediate cyclic strain amplitudes and low cyclic strain amplitudes the stress response revealed a gradual decrease of both the tensile stress (sT) and compressive stress (sC) components from the onset of fully-reversed cyclic deformation [Fig. 9]. 4.5. Mechanisms governing cyclic stress response The microscopic mechanisms, which control the variation of cyclic stress amplitude during fully-reversed strain amplitudecontrolled deformation, are dependent on the intrinsic influences of microstructure of this high strength alloy steel AerMet 100 and strain range. The plausible microscopic mechanisms governing stress response of this high strength alloy steel can be ascribed to the concurrent and mutually interactive influences of (a) Deformation in the plastic domain causes a gradual increase in the density of dislocations present in the microstructure by dislocation multiplication. (b) An interaction of the mobile dislocations with the secondphase carbide particles dispersed through the microstructure coupled with an initial interaction of the moving dislocations with both the martensite laths and the grain boundaries. (c) With time, an overall increase in the mutual interaction of a dislocation with other moving dislocations.

0.92 0.90 100

10

1

10

2

10

3

10

4

Fatigue Life (Nf) Fig. 7. Variation of normalized stress (sN/s1) with fatigue life (Nf) for alloy steel AerMets 100 when cyclically deformed at room temperature (27 1C).

The dislocation–dislocation interactions coupled with an interaction of the mobile dislocations with the intrinsic microstructural features and grain boundaries is responsible for the observed initial hardening or strengthening during fully-reversed cyclic deformation. When the local stress concentration (s*) caused by the progressive build-up of dislocations at a second-phase particle

K. Manigandan et al. / Materials Science & Engineering A 601 (2014) 29–39

35

over the mechanisms responsible for the observed initial hardening, i.e., dislocation–dislocation interactions, dislocation-microstructural feature interactions, and dislocation-second-phase carbide particle interactions, resulting in the occurrence of only softening from the onset of fully-reversed cyclic deformation. 4.6. Cyclic stress–strain response

Δ

AerMet® 100 T=27°C

T/2(%)= 0.6

Stress(MPa)

Tensile Stress

CompressionalStress

When this high strength steel is subject to cyclic strain amplitude the resultant stress amplitude will tend to change with continued cycling. The cyclic stress–strain response [CSSR], cyclic hardening and/or softening behavior is a measure of this transient response and is useful in designing and developing materials having improved cyclic fatigue resistance. Under fully-reversed, total strain amplitude-controlled cyclic straining, the resultant stresses developed in this high strength alloy steel AerMet 100 may be invariant during cycling, i.e., they may remain stable, either increase (harden) or decrease (soften) with continued cycling. For the microstructure that tends to either cyclically harden or soften, a stable stress, referred to as the saturation stress, is generally taken at specimen half-life (Nf/2). Since the chosen steel experienced softening at all values of cyclic strain amplitude tested the cyclic stress at half-life was chosen. The cyclic stress versus strain curve is used to characterize the steady state behavior and is obtained by plotting the stabilized stress with plastic strain amplitude (ΔεP/2), both taken at the first cycle and at specimen half-life, and conforms to the equation 0

Δs=2 ¼ K 0 ðΔεp =2Þn

ð3Þ 0

In this expression K is the cyclic strength coefficient (2086.4 MPa for specimen half-life) and the exponential term (n0 ¼ 0.137 for the specimen half-life curve) is the cyclic strain Fig. 9. Variation of tensile stress and compressive stress with cycles for alloy steel AerMets 100, at (a) high cyclic strain amplitude, (b) low cyclic strain amplitude.

3000

exceeds a critical value, microscopic crack initiation is favored to occur through rupture of the hard, brittle and essentially elastically deforming second-phase particle. With continued cyclic straining numerous such second-phase particles and martensite precipitates, either independently dispersed through the microstructure of alloy steel AerMet 100 or present as agglomerates, tend to fail by cracking because of the following:

2000

Stress (MPa)

AerMet® 100 T=27°C

Monotonic Cyclic (N=Nf/2) n = 0.13

n’ = 0.16

1000 900 800

With continued cyclic straining the fine microscopic cracks initiated at the interfaces of both the martensite micro-constituent and the second-phase carbide particle rapidly propagate through the microstructure and eventually link to form one or more macroscopic cracks. The initiation of numerous such fine microscopic and macroscopic cracks coupled with their concurrent growth through the microstructure of this high strength alloy steel AerMets 100 results in visible softening, i.e., a noticeable decrease in load carrying capability of the microstructure, eventually culminating in failure by fracture. At the higher cyclic strain amplitudes these concurrently occurring mechanisms dominate

K =2009 MPa K’=2086 MPa

700 600

0.01

0.1

1

Plastic Strain (%) 3000

Monotonic Cyclic(N=1)

AerMet® 100 T=27°C 2000

Stress (MPa)

(i) Their intrinsic brittleness, that is, when the local stress concentration (s*) arising from dislocation build up exceeds the strength of the second-phase particle (sparticle) or martensite precipitate (sprecipitate). (ii) Either separation or decohesion at the interface of (a) the soft, essentially ductile and plastically deforming steel matrix with the hard, brittle and elastically deforming martensite microconstituent, and (b) soft matrix with the hard second-phase carbide particle.

n = 0.13

n’ = 0.17

1000 900

K =2009 MPa K’=2204 MPa

800 700 600

0.01

0.1

1

Plastic Strain (%) Fig. 10. Comparison of the monotonic stress versus strain and cyclic stress versus strain curves for AerMets 100 at room temperature (27 1C).

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hardening exponent, which provides a measure of the resistance offered by the microstructure of this high strength steel to cyclic straining. A comparison of the monotonic stress–strain and cyclic stress–strain curves is shown in Fig. 10. At the chosen test temperature (27 1C) the cyclic stress taken at half-life at all of the cyclic strain amplitudes is marginally lower than the monotonic stress indicating the cyclic state to be softer than the monotonic state (Fig. 10a). However, at the onset of fullyreversed cyclic deformation, i.e., at the end of the first cycle, the cyclic stress is marginally higher than the monotonic stress indicating hardening during the early stages of cyclic deformation (Fig. 10b) Also, the overall degree of hardening or strengthening with increasing plastic strain, quantified by the cyclic strain hardening exponent (n0 ), is observably more for this alloy steel in the cyclic state than in the monotonic state.

4.7. Cyclic fracture behavior At the chosen test temperature, i.e., 27 1C, the cyclic fatigue fracture surfaces revealed only marginal difference in topography, over the entire range of strain amplitudes examined for this chosen high strength alloy steel AerMet 100. However, on a fine microscopic scale, the fracture surfaces were found to vary as a function of cyclic strain amplitude and resultant fatigue life (Nf). Representative fractographs of the fatigue fracture surfaces of the test samples are shown in Figs. 11–13. One sample was cyclically deformed at a high strain amplitude and resultant short fatigue life, while the other chosen sample was cyclically deformed at lower strain amplitude and resultant enhanced fatigue life.

4.7.1. Total strain amplitude of 0.9%, Nf ¼692 cycles At this cyclic strain amplitude and resultant short fatigue life (692 cycles) macroscopic fracture behavior of the chosen test sample was essentially cup and cone type (Fig. 11a). Macroscopic or low magnification, observations in the scanning electron microscope (SEM) revealed the fracture surface to comprise of features reminiscent of both brittle (fine microscopic and macroscopic cracking) and “locally” ductile mechanisms. At this high strain amplitude the region of stable fatigue crack growth was essentially small and comprised of pockets of well-defined striations reminiscent of the occurrence of microplastic deformation at the “local” level (Fig. 11b). In the region immediately prior to unstable crack growth and overload were evident traces of macroscopic cracking (Fig. 11c). The region of transition from stable to unstable crack growth was short and not distinctly evident. Also present in the transition region were fine voids of varying size, their gradual growth and eventual coalescence to form one or more fine microscopic cracks (Fig. 11d). The fine microscopic and macroscopic voids and cracks (to include both the fine microscopic and macroscopic) were surrounded by “pockets” of shallow dimples of varying size (Fig. 12a) clearly indicative of ‘local’ ductile fracture mechanisms [24,27]. The isolated pockets of flat transgranular regions, evident in the domain of early crack growth, revealed fine hairline microscopic cracks (Fig. 12b) reminiscent of the occurrence of locally brittle failure mechanisms in this region. 4.7.2. Total strain amplitude of 0.7%, Nf ¼3634 cycles At this cyclic strain amplitude and resultant enhanced fatigue life of 3634 cycles macroscopic fracture behavior of the alloy steel AerMets100 sample was essentially normal to the far-field stress

Striations

200μ

Macroscopic cracks

50μm

Voids

10μm

2μm

Fig. 11. Scanning electron micrographs of the fatigue fracture surface of AerMet 100s sample cyclically deformed at a strain amplitude of 0.095%, with a fatigue life of 92 cycles, showing: (a) Overall morphology of failure. (b) Regions of stable crack growth pockets of well-defined striations. (c) Macroscopic cracking in the transition region between stable and unstable crack growth. (d) Voids of varying size and microscopic cracks in the region of tensile overload.

K. Manigandan et al. / Materials Science & Engineering A 601 (2014) 29–39

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Void

Microscopic

Dimple 10μm

2μm

Fig. 12. Scanning electron micrographs of the fatigue fracture surface of AerMets 100 sample cyclically deformed at a strain amplitude of 0.095%, with a fatigue life of 92 cycles, showing (a) voids, microscopic cracks and dimples in the transition region. (b) Microscopic cracks surrounded by a population of dimples in the transgranular regions of early crack growth.

200μ

20μm

Dimple

5μm

2μm

Fig. 13. Scanning electron micrographs of the fatigue fracture surface of AerMets 100 sample cyclically deformed at a strain amplitude of 0.08%, with a fatigue life of 3634 cycles, showing: (a) Overall morphology of failure. (b) Region of early microscopic crack growth. (c) Region of unstable crack growth prior to overload. (d) A noticeable population of dimples covering the overload fracture surface.

axis (Fig. 13a). The transgranular pockets in the region of early microscopic crack growth when viewed at higher magnification revealed fine shallow pockets of striations intermingled with fine microscopic voids of varying size and shallow dimples (Fig. 13b). The region of unstable crack growth immediately prior to overload was microscopically rough and comprised of fine microscopic cracks intermingled with a random distribution of shallow dimples and small yet distinct pockets of near featureless transgranular regions (Fig. 13c). The region of overload revealed an observable

population of dimples of varying size (Fig. 13d) reminiscent of the “locally” occurring ductile failure mechanisms. Overall, this sample revealed a fewer number of fine microscopic cracks and macroscopic cracks when compared to the sample that was deformed at the higher cyclic strain amplitude and resultant enhanced fatigue life (Nf). Fractographic observations of the two samples deformed at two different cyclic strain amplitudes revealed that the degradation in cyclic fatigue life at the higher cyclic strain amplitude occurred as

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Dislocations

Dislocations

100n

50nm

Slip Lines

100n Fig. 14. Transmission electron micrographs of the fatigue fracture surface of AerMets 100 sample that was cyclically deformed at plastic strain amplitude of 0.095 pct., showing: (a) Non-uniform distribution of dislocations in the grains. (b) High magnification of (a) showing dislocation distribution within a single grain. (c) Formation and presence of non-uniform dispersion of very fine persistent slip bands within the grains.

a result of a drastic change in ‘microscopic’ fracture mode to predominantly brittle intergranular with pockets of ductile transgranular regions. The occurrence of plastic strain inhomogeneity at the fine microscopic level was evident from both a non-uniform distribution of dislocations through the microstructure (Fig. 14a and b) coupled with the formation and presence of well-defined persistent slip bands within the grains (Fig. 14c). The fine slip bands were observed in only few grains at the higher allowable magnifications of the SEM. With continued cyclic straining the slip bands tend to impinge upon the grain boundaries causing the “local” stress concentration (say s**) to occur at this region. When the local stress concentration exceeds the strength of the grain boundary microcrack initiation is favored to occur. The fine microscopic cracks grow and coalesce with each other to form one or more macroscopic cracks.

5. Conclusions Based on the results of an investigation aimed at understanding the strain-controlled low cycle fatigue response and resultant fracture behavior of the high strength alloy steel, AerMets 100, following are the key findings. 1. Optical microstructure is quite typical of a high strength alloy steel and revealed a combination of carbon-rich and carbondepleted regions. A higher carbon and alloy content in this steel resulted in a greater volume fraction of the micro-constituent martensite in the carbon-rich regions. The presence and overall morphology of martensite was in the form of lath.

2. This high strength alloy steel exhibited a linear trend for the variation of log elastic strain amplitude (Δεe/2) with log reversals-to-failure (2Nf), and log plastic strain amplitude (ΔεP/2) with log reversals-to-failure (2Nf) in conformance with the Basquin relationship and the Coffin–Manson relationship. 3. Cyclic stress response of this high strength alloy steel revealed a combination of initial hardening for the first few cycles followed by stability for large portion of fatigue life before culminating in rapid softening to failure at the lower cyclic strain amplitudes and intermediate cyclic strain amplitudes and resultant enhanced cyclic fatigue life. However, at the higher cyclic strain amplitudes and concomitant short fatigue life this high strength alloy steel revealed gradual softening to failure from the onset of fully-reversed cyclic deformation. 4. Fracture morphology was different at both the macroscopic and fine microscopic levels over the entire range of cyclic strain amplitudes examined. At the higher strain amplitudes macroscopic observations revealed fracture to be essentially ductile with microscopic features reminiscent of ‘locally’ ductile and brittle failure mechanisms. At the lower strain amplitudes both macroscopic and fine microscopic observations revealed fracture to be essentially ductile with features reminiscent of locally occurring ductile mechanisms. 5. The occurrence of non-uniform distribution of dislocations through the grains coupled with the formation and presence of persistent slip bands within the grains is “conducive” for ‘local’ stress concentration at the grain boundary regions and the resultant intergranular failure as evident by both fine microscopic and macroscopic cracking along the high angle grain boundaries.

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References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14]

M. Sudio, S. Hashimoto, S. Kambe, Int. ISIJ 23 (1983) 303–311. M. Sudo, T. Iwai, Int. ISIJ 23 (1983) 284–302. I. Kim, S. Rachel, W. Dahl, Steel Res. 58 (1987) 186–190. B.Y. Choi, G. Krauss, D.K. Matlock, Scr. Metall. 22 (1988) 1575–1580. D.K. Matlock, G. Krauss, Mater. Sci. Eng. A 165 (1) (1993) 1–8. G.A. Cingara, Y. Ososkov, M.K. Jain, D.S. Wilkinson, Mater. Sci. Eng. A 516 (2009) 7–15. Y. Furuya, S. Torizuka, E. akeuchi, M. Bacher Hochst, M. Kuntz, Mater. Des. 37 (2012) 515–520. D.Y. Wei, J.L. Gu, H.S. Fang, B.Z. Bai, Z.G. Yang, Int. J. Fatigue 26 (2004) 437–442. Ahmad Zare, A. Ekrami, Mater. Sci. Eng. A 530 (2011) 440–445. U. Liedl, S. Traint, E.A. Werner, Comput. Mater. Sci. 25 (2002) 122–128. A.S. Khan, R. Liang, Int. J. Plast. 15 (1999) 1089–1109. Online material manual of AerMet steels. Carpenter Technology Corporation, Reading, PA, USA, 2004. Lee, Eun. U., Fatigue Crack Growth in AerMet 100 Steel, Final Report no. NADC 91111-60, Naval Air Development Center, Warminster, PA, USA, October 1991. G.Y. Lai, W.E. Wood, R.A. Clark, V.F. Zackay, E.R. Parker, Metall. Trans. 5 (1974) 1663–1670.

[15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26]

[27]

39

G.Y. Lai, Mater. Sci. Eng. 19 (1975) 153–156. F. Zia Ebrahimi, G. Krauss, Metall. Trans. 14A (1983) 1109–1119. F. Zia Ebrahimi, G. Krauss, Acta Metall. 32 (1984) 1767–1777. B.C. Kim, S. Lee, D.Y. Lee, N.J. Kim, Metall. Trans. 22A (1991) 1889–1892. R.O. Ritchie, B. Francis, W.L. Server, Metall. Trans. 7A (No. 6) (1976) 831–838. R.O. Ritchie, R.M. Horn, Metall. Trans. 9A (1978) 331–341. B.I. Sandor, Fundamentals of Cyclic Stress and Strain, University of Wisconsin Press, Madison, Wisconsin, USA, 1972. L.F. Coffin , Trans. ASM 76 (1954) 931–940. S.S. Manson: NASA Technical Note, National Aeronautics and Space Administration, 1954, pp. 2933–2940. E.A. Starke , G. Lutjering, in: M. Meshii (Ed.), Fatigue and Microstructure, American Society for Metals, Metals Park, Ohio, USA, 1979, pp. 205–240. K. Manigandan, T.S. Srivatsan, T. Quick, S. Sastry, M.L. Schmidt, Mater. Des., 2014 (in press). ASTM E8-2013: Standard Test method for Tension Testing of Metallic Materials, American Society for Testing and Materials, ASTM, Race Street, Philadelphia, PA, 2013. W.E. Wood, Eng. Fract. Mech. 7 (1975) 219–228.