Materials and Design 56 (2014) 146–157

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Materials and Design

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Low cycle fatigue behavior of Cr–Mo–V low alloy steel used for railwaybrake discs

0261-3069/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.matdes.2013.10.093

⇑ Corresponding author. Tel./fax: +86 10 51683300.E-mail address: jmhan@bjtu.edu.cn (J. Han).

Zhiqiang Li, Jianmin Han ⇑, Weijing Li, Like PanSchool of Mechanical, Electronic and Control Engineering, Beijing Jiaotong University, Beijing 100044, PR China

a r t i c l e i n f o

Article history:Received 20 August 2013Accepted 25 October 2013Available online 12 November 2013

Keywords:Low cycle fatigueCr–Mo–V steelFatigue behaviorElevated temperatureMetallographyFractography

a b s t r a c t

The cyclic stress–strain response and the low cycle fatigue (LCF) behavior of Cr–Mo–V low alloy steelwhich was used for forged railway brake discs was studied. Tensile strength and LCF properties wereexamined over a range from room temperature (RT) to 600 �C using specimens cut from circumferentialdirection of a forged disk. The fully reversed strain-controlled LCF tests were conducted at a constant totalstrain rate with different axial strain amplitude levels. The cyclic strain–stress relationships and thestrain–life relationships were obtained through the test results, and related LCF parameters of the steelwere calculated. The studied steel exhibits cyclic softening behavior and behaves Masing type, especiallyat higher strain amplitudes. At higher than 600 �C, carbide particles aggregated and a decarburized layerdeveloped near the specimen surface. Micro voids distribute within the depth of 50 lm from thespecimen surface could coalesce with fatigue cracks. Multiple crack initiation sites were observed onthe fracture surface. The oxide film that generated at 600 �C covered the fatigue striations and acceleratedthe crack propagation. Final fracture area with bigger and deeper dimples showed better ductility athigher temperature. The investigated LCF behavior can provide reference for brake disc life assessmentand fracture mechanisms analysis.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Cr–Mo–V low alloy steels which have high hardenability, favor-able strength and toughness are widely used material for manyapplications that include the mechanical, transportation and petro-leum industries. Some of the reasons for this are that the relativelylow percent volume of alloying elements make the steel more eco-nomical and the toughness and ductility can be further improvedby keeping low impurities content through the use of electroslagrefining [1,2]. With appropriate heat treatment, such as quenchingand tempering (Q&T), it is possible to achieve a high level ofstrength and toughness along with a high degree of heat-resis-tance. A further advantage of Cr–Mo–V low alloy steel is that ithas relatively high heat conductivity, low thermal expansion coef-ficient and good heat shock resistance [3]. This enables the mate-rial to operate at high temperatures for long durations and makeit an ideal material for railway brake discs [4–6].

In railway transportation industries, brake discs are essentialsafety components, especially for high speed trains. During brak-ing, the brake pads press against the friction surface of brake discsand transform the kinetic energy of the railway vehicles into heat.The friction heat generated during braking causes significant

temperature variation on brake discs. In a typical railcar, the peaktemperature on the surface of brake disc can reach in excess of600 �C. As the trend in transportation is towards an increasing ofmaximum speed and braking loads, the higher level of thermaland friction loadings during braking can induce thermal stressand local plastic strains near the friction surface [7–11]. Becauseof repeated heating and cooling caused by braking process, severalissues are of concern. The first is that the thermal fatigue cracksmay appear under the action of the complex thermal–mechanicalload. The heat shock and crack propagation of brake discs becomecommon issues for railway transportation field these years[8–10,12]. The thermal fatigue of brake discs, which is essentiallya low cycle fatigue (LCF) progress, may lead potential danger tothe brake discs. However, high fatigue resistance and LCF behaviorof brake disc materials is also of great significance and it mayessentially decide the lifespan.

There have been large amount of studies that have been con-ducted to investigate the friction properties of the braking pairs.Findik [13–15] investigated the friction behavior of brake liningsmade of various ingredients. The effect of the friction environment,pressure and temperature on the pads was discussed. Furthermore,LCF behavior of different kind of brake disc materials is also a hottopic. Since 1997, Samrout and El Abdi [5,6] have started re-searches on the fatigue behavior of 28CrMoV5-8 brake disc steel,which had been used in French high speed trains for many years.

Z. Li et al. / Materials and Design 56 (2014) 146–157 147

A series of isothermal and anisothermal experiment were con-ducted at strain range 1.4% under cyclic uniaxial tensile and com-pressive load. An anisothermal fatigue damage model based on theManson–Coffin relationship was presented, as well. Singh et al.[16] evaluated the effect of Q&T heat treatment on the LCF behav-ior of a traditional French steel 15CrMoV6 at RT. The fatigue resis-tance of the steel after Q&T process was lower than as-receivedcondition because of a reduction of the fracture toughness. Šamecet al. [17] conducted a series of monotonic test and strain-con-trolled LCF experiments on studying the cast iron EN-GJS-500-7used for brake disc material and the LCF behavior was investigatedfrom RT to 400 �C. The basic LCF parameters of the cast iron wereproposed in his work. LCF behavior of Cr–Mo–V steels used forother industry applications was also investigated. Zhang et al.[18] conducted a series of LCF tests at room temperature to inves-tigate the LCF behavior of a Cr–Mo–V high-speed steel, which isused for cold forging tool. The microstructure of the steel wasinvestigated and the brittle phases were found to be the sourceof the fatigue cracks. The cracks could initiate from both the spec-imen surface and the interior area. Kneifl et al. [19] proposed amodel describing the life of Cr–Mo–V and Cr–Ni–Mo–V steels atcombination of LCF and creep and elaborated a computationalmethod to calculation the fatigue life. Luo et al. [20] evaluatedthe LCF life of a micro-alloyed high strength steel using energy-based prediction method. Plastic strain energy per cycle was con-sidered to be an important parameter for lifespan evaluation.

The lifespan and operational safety issues of brake discs are ofgreat significance. Brake discs are subject to severe temperaturevariations during its lifetime, the LCF behavior of brake disc steelat different temperature levels must be established prior to itsapplication in the railway vehicles. However, there has only afew research conducted on LCF behavior of railway brake discsteels. In this study, a kind of indigenously-developed Cr–Mo–Vlow alloy steel for railway brake discs was used as an examplefor LCF behavior investigation. The studied steel, having composi-tion similar to 28CrMoV5-8 steel but with higher carbon content,has been developed following the electroslag refining (ESR) processand kept in low impurity contents.

The aim of this study is focused on the LCF behavior and frac-ture mechanism of the brake disc steel. The LCF tests were con-ducted under strain-controlled condition at RT and elevatedtemperatures. Basing on the Manson–Coffin–Basquin and Ram-berg–Osgood equations, the strain–life relationship and cyclicstress–strain behavior of the material were all obtained by the testresults. Moreover, in order to get a better understanding of crackinitiation and propagation mechanism of the studied steel duringLCF process, metallographic and fractographic observations arenecessary. The related methods and test results in this study canprovide basic reference for life assessment of brake discs basingon the elastic and plastic strain methods. The micro-mechanismof crack initiation and propagation can be used for comparisonwith type of failure of brake discs, as well as other componentsworking with frequently temperature variation.

Fig. 1. Geometry of the monotonic tensile test specimen (unit: mm).

2. Experimental procedure

In this study, LCF properties of a low alloyed brake disc steelwere studied. The steel contains 0.32% C (weight percent) andother chemical composition is close to 28CrMoV5-8 brake discsteel. The material was produced by electroslag remelting processand the gas and impurities content of the steel was strictly con-trolled within a low level. The casting ingot was firstly forged intoa disc having a 615 mm diameter and 55 mm thickness. The initialforging temperature was 1150 �C and ended at 880 �C. Then smallbars were cut along the circumferential direction from the forged

disk and then machined into monotonic tensile and LCF tests spec-imens. After that, a Q&T heat treatment process, which included oilquenching after austenitizing at 880 �C for 1 h and air cooling aftertempering at 660 �C for 2 h, was applied to the specimens beforefinishing machining.

Prior to performing the LCF tests, monotonic tensile tests atdifferent temperatures were conducted to evaluate the mechanicalproperties of the material. Tensile tests were conducted using anelectronic universal testing machine CMT5105 with a load capacityof 100 kN that was equipped with a resistance heating furnace. Thetensile tests were performed at room temperature (RT), 200 �C,400 �C, 600 �C, and 800 �C with three replications at each temper-ature level. All tensile test specimens (Fig. 1) were designedand machined in accordance with the Chinese National StandardGB/T 4338-2006 ‘‘Metallic materials – Tensile testing at elevatedtemperature’’ [21]. The working section of the cylindrical speci-mens had a diameter of 4 mm. Strain was measured continuouslywith an extensometer of gauge length of 20 mm and travel lengthof 10%.

The LCF tests were carried out using a 250 kN capacity servo-hydraulic machine (MTS-810) equipped with an electric resistanceheating furnace. The LCF test specimens (Fig. 2) were machinedaccording to Chinese National Standard GB/T 15248-2008 ‘‘TheTest Method for Axial Loading Constant Amplitude Low-Cycle Fati-gue of Metallic Materials’’ [22]. The specimens with a diameter of6.35 mm and gage length of 25 mm were finely polished beforetesting to ensure a consistent surface finish. Temperature calibra-tion was conducted before the fatigue tests. Three thermocoupleswere fixed on the top, middle and bottom sections of a specimen,respectively. These thermocouples were used to measure surfacetemperatures of the specimen to ensure an uniform temperaturedistribution. The temperatures were controlled in a constant valueand temperature differences between the specimen and the targettemperature were strictly controlled within ±2 �C.

The strain controlled LCF tests were carried out at RT, 200 �C,400 �C and 600 �C, respectively. The total axial strain amplitudeswere controlled at different levels of ±1.0%, ±0.8%, ±0.6%, ±0.4%,and a lower level from 0.25% to 0.35%. Two replications were testedat each strain amplitude with a fully push–pull mode (minimum tomaximum strain ratio, R = �1). A triangular waveform signal with aconstant strain rate of 5.0 � 10�3 s�1 was used. The longitudinalstrain was measured continuously with an extensometer duringeach of the LCF test. The cyclic stress–strain data was recordedand the hysteresis loops were obtained. The plastic strain rangewas measured from the width of the monitored hysteresis loop.The stress ranges were examined by the peak and valley stress val-ues during each strain cycle which were determined from the tipsof the hysteresis loop. The number of cycles which leads to a dropof 25% of the tensile stress was taken as fatigue life and the mid-lifecycle was taken as stable hysteresis loop.

After the tests were completed, selected specimens wereexamined by scanning electron microscope (SEM) on theirmicrostructures and fracture features. The thin sections for metal-lographic analysis were cut off from the monotonic test specimens

Fig. 2. Geometry of the LCF test specimen (unit: mm).

148 Z. Li et al. / Materials and Design 56 (2014) 146–157

in a transverse direction near the fracture surface. The other sam-ples are longitudinally cut off from LCF test specimens in order toanalyze the microstructures and the propagation path of fatiguecracks. These specimens were polished and etched with a nitaletchant (4% nitric acid 96% ethanol) and then examined by a FEIQuanta 400 SEM in the secondary electron mode. The fractographicexaminations of the fracture surface of LCF test specimens wereconducted on a Hitachi S-530 SEM.

3. Results and discussions

3.1. Mechanical properties and microstructure

Mechanical properties of the steel obtained by tensile tests fromRT to 800 �C are shown in Fig. 3. During each test, Young’s modulus(E), ultimate tensile strength (UTS), yield point (YP), as well aselongation were characterized. It can be seen that with the in-crease in temperature, all mechanical properties of the steel de-creased except elongation. There was only a slight reduction inUTS, YP and E until 500 �C while the elongation was essentiallythe same. When the temperature was above 600 �C, the UTS, YPand E demonstrated a significant decline and the elongation in-creased abruptly. At 800 �C, the UTS and YP was only 10% of thatat RT, while the E was 38% of that at RT. And there was a sharp risein average elongation which was more than 150%. Basically, thestudied material exhibited a little higher yield strength, Young’smodulus and elongation than a typical 28CrMoV5-8 at RT [4,5].

The micrograph in Fig. 4a shows that the microstructure of thestudied steel primarily consists of tempered martensite at RT.Large amount of finely dispersed carbides distribute in ferrite sub-strate and the some of the grains maintain the same origin grainorientation of the untempered martensite. Fig. 4b and c show themicrostructure of the specimens after tensile tests that wereperformed at 600 �C and 800 �C. The microstructure of the mono-tonic tensile specimen was similar below 600 �C and showed good

Fig. 3. Mechanical properties at RT and elevated temperatures.

Fig. 4. The microstructure of the steel after monotonic tensile test at (a) RT, (b)600 �C and (c) 800 �C.

stability. After testing at 800 �C, the tiny spherulitic carbidesdistinctly aggregated and became larger. The carbides near the sur-face reduced and the ferrite grain size increased significantly. Itindicates that severe decarburization and grain growth occurred

Table 1The plastic strain ranges (Dep) at RT and elevated temperatures.

Strain amplitude (Det=2) ±1.0% ±0.8% ±0.6% ±0.4%

Dep (RT) 0.01014 0.00670 0.00320 0.00036Dep (200 �C) 0.01076 0.00740 0.00412 0.00124

Z. Li et al. / Materials and Design 56 (2014) 146–157 149

near the specimen surface at higher temperature. The change ofmicrostructure such as spheroidized carbides particles and graingrowth should be taken into consideration when operating condi-tions of the material are around 800 �C which is higher than theAc1 temperature of the steel.

Dep (400 �C) 0.01158 0.00760 0.00420 0.00106Dep (600 �C) 0.01322 0.00906 0.00554 0.00180

3.2. Cyclic stress–strain behavior

During the LCF tests, the cyclic stress and strain produce dis-tinct hysteresis loops. In contrast to monotonic testing, the cyclicloading does not lead to a unique relationship between stressand strain, but rather to a hysteresis loop for each cycle, whichshows the relationship between the total strain range (Det) andthe total stress range (Drt). In this study, the hysteresis loop athalf–life was taken as the stable loop. The stable hysteresis loopsat different cyclic strain range and temperature levels were plottedin Fig. 5. The plastic strain range (Dep) was measured from thewidth of the stable hysteresis loop, which is shown in Table 1.For the tests at same strain range, Dep increased with testing tem-perature. It attributes to the decrease in YP with the increase intemperature. At low strain range, the plastic strain also becomeslower and the hysteresis loop tends to become a straight line(Fig. 5a and d). It correlates well with the previous works[17,18]. This phenomenon of ferrous materials represents thatthere is only cyclic elastic deformation occurred in the materialand the slope of the straight line approximately equals to theYoung’s modulus E.

By connecting the vertices of the hysteresis loops, the cyclicstress–strain curves can be obtained. It shows the stable stress–strain behavior of the material during LCF test. The relationship

Fig. 5. Stable stress–strain hysteresis loops at (a

between stress amplitude Drt/2 and strain amplitude Det/2 canbe described by the cyclic stress–strain equation:

Det

2¼ Dee

2þ Dep

2¼ Drt

2Eþ Drt

2K 0

� �1=n0

ð1Þ

where Dee/2 is the elastic strain amplitude, Dep/2 is the plasticstrain amplitude, K 0 is the cyclic strength coefficient, and n0 is thecyclic strain-hardening exponent. Eq. (1) is in same form with theRamberg–Osgood equation which is used for describing the mono-tonic stress–strain relationship. The plastic term of Eq. (1) can alsobe written as:

Drt=2 ¼ K 0ðDep=2Þn0

ð2Þ

The plot of Drt/2 versus Dep/2 in Eq. (2) shows a linear fit rela-tionship when plotted in logarithmic coordinates. Basing on Eq. (2),the cyclic stress–strain curves at different temperatures are shownin Fig. 6. The cyclic strength coefficient K 0 and the cyclic strain-hardening exponent n0 for cyclic conditions were calculated usinga best fit linear regression method. The parameters at differenttemperatures are shown in Table 2.

) RT, (b) 200 �C, (c) 400 �C, and (d) 600 �C.

Fig. 6. Cyclic stress–strain behavior at RT and elevated temperatures.

Fig. 7. Cyclic stress response during LCF tests at (a) RT and (b) 600 �C.

150 Z. Li et al. / Materials and Design 56 (2014) 146–157

3.3. Cyclic stress response

In the strain-controlled LCF test, the cyclic stress response is animportant feature of a material as it represents the path of thestress amplitude by which the material arrive at their final stresslevel. With the peak and valley true stresses which were recordedduring each tensile and compressive cycle, the cyclic stress re-sponse curves at different strain amplitudes were plotted and areshown in Fig. 7.

It is obvious that the cyclic stress response of the material is clo-sely related to the total strain ranges (Det). The fatigue life in-creased obviously when the strain levels decreased at eachtemperature level. Under all the test conditions, the stress ampli-tude decreased with the increase in the number of cycles, whichmeans the steel exhibits cyclic softening behavior. The decreasewas greater and the softening behavior becomes more pronouncedat higher strain amplitude levels. For example, in the first cycle atRT (Fig. 7a), the stress amplitude of the test at strain amplitude of1.0% was 1050 MPa while the stress amplitude at a strain ampli-tude of 0.4% was 832 MPa. The stress amplitude of the two speci-mens reduced by 145 MPa and 56 MPa before the failure of thespecimens and the decreasing rate of stress amplitudes were13.8% and 6.7%, respectively. Under the lowest total strain range,such as strain amplitude of 0.33% at RT, the stress response exhib-ited more stability. The cyclic stress responses at other tempera-ture levels showed similar behaviors to those at RT. However, athigher temperature levels, there were great decrease in stressamplitudes and slight decrease in fatigue lives, which is shownin Fig. 7b. Moreover, the cyclic stress response of testing at ampli-tudes above 0.4% showed same trend at 600 �C.

Similar to other steels with relatively high strength, the studiedsteel behaved cyclic softening under cyclic load. There were alsothree different stages of the cyclic stress responses at all strainranges and temperature levels [18]. At the beginning a slight neg-ative slope of stress amplitude occurred and it showed a cyclicsoftening stage of the material. This phenomenon can be seen more

Table 2The cyclic strength coefficient and the cyclic strain-hardening exponent at RT andelevated temperatures.

Temperature (�C) K 0 (MPa) n0

RT 1175 0.0488200 1345 0.0906400 1036 0.0603600 712 0.0536

clearly in Cartesian coordinates than logarithmic coordinates. Thena following smooth and less evident descending stage contributedthe main LCF life of the material. In this stage, the stress responseexperienced a less decrease, and stress amplitude was nearly con-stant at low strain amplitudes. In the third region of the cyclicstress response curve, there was a sharp descending stage of thestress amplitude. It represents the instantaneous failure of thematerial with the rapid propagation of fatigue cracks.

3.4. Cyclic strain–life relationships

Fig. 8 shows the relationship between true strain amplitude andthe reversals to failure. The fitted curves were plotted in logarith-mic coordinates. 2Nf is the number of strain reversals to failure foreach tested specimen, which is the double of the number of cyclesto fatigue failure (Nf). The relationship between the elastic strainamplitude (Dee/2) and 2Nf can be described by the Basquin [23]equation:

Dee

2¼

r0fEð2Nf Þb ð3Þ

The relationship between the plastic strain amplitude (Dep=2)and 2Nf can be described by the Manson [24]–Coffin [25] equation:

Dep

2¼ e0f ð2Nf Þc ð4Þ

The total true strain amplitude (Det=2) is the sum of the elasticstrain amplitude and the plastic strain amplitude. Therefore, thetotal strain amplitude-number of reversals to failure(Det=2� 2Nf ) is expressed as:

Det

2¼ Dee

2þ Dep

2¼

r0fEð2Nf Þb þ e0f ð2Nf Þc ð5Þ

where r0f is the fatigue strength coefficient, b is the fatigue strengthexponent, e0f is the fatigue ductility coefficient, and c is thefatigue ductility exponent. In double logarithmic coordinates, the

Fig. 8. Strain-life curves at (a) RT, (b) 200 �C, (c) 400 �C, and (d) 600 �C.

Z. Li et al. / Materials and Design 56 (2014) 146–157 151

strain–life curves can be fitted with Eqs. (3) and (4) and plotted bytwo straight lines with slopes b and c. Intercepts with the verticalaxis of the two straight lines fitted with Eqs. (3) and (4) representsthe constants r0f =E and e0f . The Det=2� 2Nf curves were simply fittedwith Eq. (5). The LCF parameters of the material basing on Eqs. (3)–(5) are listed in Table 3. These calculated basic LCF constants corre-lates well with those measured in large majority of low alloy steels[26].

It is evident in Fig. 8 that the plastic strain is lower than elasticstrain under most strain amplitudes. At higher testing strain levels,especially at elevated temperatures, the plastic strain is higherthan the elastic strain. Otherwise, the total strain amplitude mainlyconsists of the elastic strain. It means that the plastic deformationtends to play a dominant role at temperatures around 600 �C withhigh strain levels.

When the magnitude of Dep/2 equals to the magnitude of Dee/2,the transition from LCF to high cycle fatigue (HCF) occurred. Theintersection of the elastic and plastic straight lines represents thetransition fatigue life in reversals (2Nt). By combining the Eqs. (3)and (4), the 2Nt can be expressed as:

Table 3LCF parameters of the material at RT and elevated temperatures.

Temperature (�C) r0f (MPa) b e0f c

RT 1241 �0.0684 2.5573 �0.9657200 1193 �0.0751 0.2660 �0.6312400 994 �0.0533 0.4527 �0.7360600 705 �0.0474 0.5898 �0.7714

2Nt ¼ ðe0f E=r0f Þ1=ðb�cÞ ð6Þ

The region to the left of this point, where 2Nf is lower than 2Nt,is mainly dominated by the plastic strain and represents the LCFregion. The region to the right, where 2Nf is higher than 2Nt, isdominated by the elastic strain and represents the HCF region.The 2Nt and corresponding Det/2 at RT and elevated temperaturesare listed in Table 4. The result of 2Nt obtained using Eq. (6) showsno apparent correlation with the testing temperature levels, whilea decrease of the strain amplitude at the transition life can be ob-served by elevating the testing temperature.

The relationship between number of reversals to failure (2Nf)and testing temperature levels are shown in Fig. 9. It is obviousthat testing at higher strain amplitude or temperature levels leadsto shorter fatigue lives. When testing at strain amplitude ±1.0%, thefatigue lives are nearly the same at different temperature levels.The LCF lives at strain amplitude ±0.4% decrease obviously withthe increase in testing temperature. The LCF lives of the studiedsteel showed same trend with traditional brake disc steel28CrMoV5-8 which was tested at strain amplitude ±0.7% by El Abdi

Table 4The transition fatigue life from LCF to HCF at RT and elevated temperatures.

Temperature (�C) 2Nt Det/2

RT 685 0.00934200 819 0.00870400 591 0.00413600 840 0.00327

Fig. 9. Number of reversals to failure versus temperature at different strainamplitudes.

152 Z. Li et al. / Materials and Design 56 (2014) 146–157

[27]. Furthermore, the refined steel in present study exhibits long-er fatigue life at above 500 �C, which might attribute to its lowerimpurities content and finer grain size.

3.5. Masing-type behavior

Masing cyclic stress–strain behavior is an important propertyfor describing the hysteresis loops obtained by LCF tests. In orderto check whether a material exhibits Masing behavior or not, the

Fig. 10. The stress–strain hysteresis loops starting from a com

lower tips of stable hysteresis loops corresponding to differentstrain amplitudes of the LCF tests should be transferred to acommon origin. For a Masing-type material, the ascending stagesof all these transformed hysteresis loops follow a same track whichis called Masing curve. On the contrary, a material exhibitsnon-Masing-type behavior when the higher tips of the hysteresisloops do not match with the track. The Masing curve can bedescribed by a transformation of Eq. (1) and expressed as follow:

Det ¼Drt

Eþ 2

Drt

2K 0

� �1=n0

ð7Þ

It shows that for a Masing-type material, the Masing curve canbe described by the cyclic stress–strain curve which is magnifiedby a factor 2 [28]. The cyclic plastic strain energy per cycle(DWp) of a Masing-type material can be calculated using the fol-lowing the expression [29]:

DWp ¼1� n0

1þ n0

� �DrtDep ð8Þ

By matching all the lower tips of the hysteresis loops at half-life(stable loops shown in Fig. 5), Fig. 10 plotted the transferred hys-teresis loops at different strain amplitudes with a common originat different testing temperatures. Masing behavior is more evidentin the LCF test at higher than 200 �C (Fig. 10b–d) while the ascend-ing sides of the loops do not match up with each other perfectlyunder low strain amplitude at RT (Fig. 10a) and 400 �C (Fig. 10c).Although the investigated steel grade does not exhibit an idealMasing-type behavior concerning the hysteresis loops at lowerstrain amplitude, most of the other transferred hysteresis loops

mon origin at (a) RT, (b) 200 �C, (c) 400 �C, and (d) 600 �C.

Fig. 11. Microstructure of the steel after LCF test at (a) RT (Specimen: No. 1,Det=2 ¼ �1:0%, 2Nf ¼ 408). (b) 600 �C (Specimen: No. 16, Det=2 ¼ �1:0% ,2Nf ¼ 350). (c) 600 �C (Specimen: No. 19, Det=2 ¼ �0:4%, 2Nf ¼ 6330).

Fig. 12. The voids nucleation (a) at grain boundary and (b) around the fatigue cracktip.

Z. Li et al. / Materials and Design 56 (2014) 146–157 153

of the steel did not violate the Masing rules significantly. Inaddition, for a non-Masing-type material, Jhansale and Topper[30] observed that the nonlinear portion of the hysteresis loops

did not change appreciably for a wide range of structural metals.The upper branches of the hysteresis loops can be matchedthrough translate the hysteresis loops along the linear portion.For the steel in this study, no evident plastic strain was observedbelow strain amplitude 0.4% at RT and 400 �C. Since the DWp forlife assessment is mainly dominated by the plastic strain, the devi-ation from Masing-type description (Eq. (7)) can only be signifi-cantly influenced by the mismatch of the hysteresis loops withhigher plastic strain range. For the studied material, the cyclic en-ergy calculation and life prediction models basing on the assump-tion of Masing behavior can be used for the studied material[18,31,32].

3.6. Metallography

After the LCF test, fractured specimens were cut along longitu-dinal direction for observing the microstructure, initiation sitesand propagation path of cracks using SEM (FEI Quanta 400).Fig. 11 shows the microstructures of the specimens after testingat RT and 600 �C with strain amplitude ±1.0% and ±0.4%. It can beseen that there was no significant microstructure change whenthe strain amplitude was controlled as 1.0% (Fig. 11a and b). Themicrostructure of specimen No. 1 and No. 16 both consist of tem-

Fig. 13. The fatigue crack initiation and propagation: (a) microcracks nucleationfrom the specimen surface, (b) secondary cracks during propagation.

154 Z. Li et al. / Materials and Design 56 (2014) 146–157

pered martensite. Unlike the microstructure shown in Fig. 11a andb, there was a thin decarburized layer near the surface of specimenNo. 19 (strain amplitude 0.4% in Fig. 11c). The occurrence of thedecarburized layer on specimen No. 19 results from working inhigh temperature atmosphere for long period of time which wasabout 170 min. The 2Nf of specimen No. 19 was 17 times longerthan specimen No. 1 and No. 16, approximately. The arrows inFig. 11b and c show several visible cavities beneath the surfaceof specimens tested at 600 �C while no void was observed in spec-imen No. 1. The voids which were visible in this magnificationmostly distributed within the depth of 50 lm from the specimensurface.

Since the yield point of the studied material dropped about 40%at 600 �C (Fig. 3), dislocation can slip more easily and plastic defor-mation is more evident when testing at high temperature levels.With the plastic deformation, the voids can form on grain bound-aries or around inclusions [33]. For the studied material, only asmall amount of inclusions and impurities were found so theseare unlikely to play an important role in the void nucleation pro-cess. It can be seen in Fig. 12a that the voids nuclear at the grainboundaries of ferrite matrix which contains different content ofcarbide particles, as well as the grain boundaries between thegrains that remain the martensite orientation with large amount

of carbides. When the plastic deformation occurs in the material,the ferrite is more likely to deform as a soft phase while the local-ized non-uniform distribution of carbides causes stress concentra-tion and deformation mismatch [32,33]. Furthermore, voids werealso likely to nucleate around the fatigue crack tips, which is obvi-ous in Fig. 12b. It is evident that the orientation of the microstruc-ture near the crack tip was same to the action of plasticdeformation. During the cyclic tension and compression process,the voids grew up and finally coalesced with the crack tip, whichcan be found in Fig. 13a. The coalescence of voids and the cracktip is a typical way of the fatigue crack propagation. It leads toabrupt propagation of cracks especially when the cracks mergewith voids in large size.

There are large amount of microcracks initiate from the speci-men surface when testing above 400 �C (white arrows inFig. 13a). In contrast, there was only one main crack observed onthe surface of specimens tested at RT. When conducting the LCFtests at elevated temperatures, specimens need to be preheatedin the heating furnace for a period of time after clamped on thehydraulic machine. Oxides generated and covered on the specimensurface during the preheating process and the subsequent fatiguetests. The brittle oxide film causes stress concentration. The oxidefilm fractures more easily under the action of cyclic loads and theseinitial cracks could propagate inward. On the other hand, whendislocations slipped onto the surface continuously, the fatiguecracks could also initiate around these local sites beneath the oxidefilm [26]. Once these microcracks formed on the surface, the crackspropagated straightly inward. This way of crack propagation is dif-ferent from that in the previous study [6]. There were few crackswhich propagate in the maximum shear plane. Moreover, somecracks propagated inwards with the occurrence of either deep orshallow secondary cracks (Fig. 13b). The secondary cracks whichpropagated along 45� angle from the main crack did not extend along distance. This kind of propagation reflects the microstructureof tempered martensite with diffusive carbide particles and finegrain size is capable to maintain the direction of crack propagationand avoid the occurrence of too many crack branches. It also makesthe studied steel perform a steady propagation stage during the fa-tigue life.

3.7. Fractography

Fracture surfaces of all the LCF test specimens were examinedusing SEM (Hitachi S-530). The aim of this analysis is to distinguishdetailed characteristics of the morphology of fracture surfaces andthe failure mechanisms that occurred at different temperatures.

Figs. 14 and 15 show representative micrographs of the fracturesurfaces features of the specimens cycled to failure at 400 �C and600 �C under strain amplitude ±0.4%, respectively. In Fig. 14, fati-gue cracks initiated from the external surface of the specimensand propagated in the direction perpendicular to the principalstress axis. There were distinct regions of fatigue crack propagationshown as the appearance of alternating ridges or strips of steps onthe facture surface and multiple crack initiation sites can be ob-served near the border of the fracture surface (arrows inFig. 14a). Under higher magnification, the crack nucleation regionrevealed that there was local deformation on the surface of speci-men under the cyclic load (Fig. 14b). The local deformation type ofnucleation was described by Cottrell and Hull by using an intru-sion–extrusion pair model. Paired dislocation may cause pile-upsat the surface of the steel or against obstacles and slip bands even-tually become cracks [31]. For the studied steel, oxides generatedon specimen surface at elevated temperatures could be anotherfactor of crack initiation. In Fig. 14c, fatigue striations in the steadyfatigue growth region and some parallel secondary cracks wereclearly visible. The average striation space was about 2–3 lm at

Fig. 14. SEM fractograph of LCF specimen tested at 400 �C and Det=2 ¼ �0:4%, 2Nf ¼ 8568; (a) crack origin sites on the surface, (b) higher magnification view of the fatiguecrack initiation area ‘‘A’’, (c) fatigue striations in area ‘‘B’’ (d) final rupture area with shallow dimples.

Fig. 15. SEM fractograph of LCF specimen tested at 600 �C and Det=2 ¼ �0:4%, 2Nf ¼ 6330, (a) crack origin sites on the surface, (b) higher magnification view of the fatiguecrack initiation area ‘‘C’’, (c) fatigue striations in area ‘‘D’’, (d) final rupture area with deep dimples.

Z. Li et al. / Materials and Design 56 (2014) 146–157 155

400 �C. The fatigue striations, which reflect the ductility ofmaterials and fatigue crack propagation process, are influencedby chemical composition, microstructures, and cyclic loadings[6,34]. The fracture of LCF specimens showed transgranularmechanism during crack growth. Similarly, all of the other LCFtests at temperature no higher than 400 �C showed same failuremechanisms.

At 600 �C, multiple crack initiation sites can be observed aswell(arrows in Fig. 15a). In Fig. 15b, oxide film on the fracture

surface is clearly visible because the specimens were tested in anair atmosphere and the fracture surface was exposed in high tem-perature for long time. Although some of the fatigue features arecovered by the oxide film, fatigue striations still can be observedbeneath the oxide film on the surface (Fig. 15b and c). Moreover,the oxide film generated on the surface of fatigue cracks is one ofthe factors which tend to accelerate the crack propagation, whichmainly lead to the reduction of fatigue life with the increase intemperature. It is reflected in Fig. 9 that a larger reduction of 2Nf

Fig. 16. SEM graph of the voids beneath the specimen surface at Det=2 ¼ �1:0%, (a)400 �C and (b) 600 �C.

156 Z. Li et al. / Materials and Design 56 (2014) 146–157

occurred at low strain amplitudes and higher temperatures, whichattribute to a longer period of testing time. This is consistent withthe previous work [35].

The final rupture area of the LCF specimens is shown in Fig. 14dand 15d. The size of dimples reflects ductility when the specimensfinally ruptured. The dimples after testing at 400 �C were shallowand small while dimples at 600 �C were much bigger and deeper.There were also carbide precipitate particles in the big dimplesat 600 �C. It indicates a better ductility of the material, and thespherulitic carbide particles tend to be aggregated and became lar-ger at higher temperatures.

Fig. 16 shows the crack initiation sites with voids beneath thespecimen surface. Under higher strain amplitudes, large plasticdeformation near the specimen surface can induce the voids nucle-ation and growth. Fatigue cracks initiated and propagated alongwith the formation of series of cavities. Similar to the voids distrib-uted in the longitudinal section in Fig. 11, the voids on the fracturesurface mostly distributed within a depth of 50 lm from the spec-imen surface. The nucleation of voids starts at second phase parti-cles, inclusions, or grain boundary of the material [31,33,36]. Thiskind of distribution of voids within a certain depth makes small fa-tigue cracks propagate inward faster in the early stage. When thecracks arrived a certain distance from the specimen surface, therewas coalescence of cracks and voids.

4. Conclusions

In this paper, LCF behavior of Cr–Mo–V steel developed forbrake discs was evaluated using strain-controlled uniaxial loading

tests. The tests were conducted at RT and elevated temperatures.The chemical composition of the studied steel is close to28CrMoV5-8 with higher carbon content and lower impurities.The basic LCF parameters were obtained from the tests and pre-sented in this article. The following conclusions of this work wereobtained:

(1) The mechanical properties and microstructure of this steel isstable until 600 �C. At higher temperature, the changes ofmicrostructure must be taken into account when determin-ing mechanical performance or predicting the LCF life.

(2) Hysteresis loops of the LCF tests were observed at strainamplitudes of ±1.0%, ±0.8%, ±0.6%, ±0.4% and lower level.The LCF behavior was analyzed at these levels, including cyc-lic stress–strain behavior, stress responses, strain–life rela-tionship and Masing analysis. Basic LCF parameters atdifferent testing temperatures were obtained.

(3) Stress responses at different strain ranges and temperaturelevels exhibited similar cyclic softening behavior for thismaterial. The softening behavior was more evident at highstrain levels. By transferring the stable hysteresis loops intoa same origin, the ascending sides of the loops did not vio-late the Masing criterion significantly. The material andbehaves Masing-type approximately and the Masing modelcan be used in analyzing the studied material.

(4) The metallographic analysis showed that carbide particlesaggregated and decarburized layer became obvious at tem-peratures higher than 600 �C. Micro voids nucleated at grainboundaries because of the deformation mismatch during thecyclic tension and compression process. Microcracks initi-ated from the surface of specimens more easily at higherthan 400 �C. The coalescence of fatigue cracks and the voidswithin the depth of 50 lm from the specimen surface lead toabrupt propagation of cracks. There were not long secondarycracks observed during the crack propagation.

(5) Fractographic observations showed multiple crack initiationsites and distinct fatigue striations on the fracture surfaces.Oxide film that generated at 600 �C covered on the fatiguestriations and accelerated the crack propagation. Final frac-ture zone of the specimens showed predominantly ductilefracture mechanism and the micro dimples with carbideparticles on the bottom were bigger and deeper at highertemperature.

Acknowledgement

The authors would like to acknowledge the support of the Na-tional Natural Science Foundation of China (Grant No. 51271014).

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