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Journal of Materials Engineering andPerformance ISSN 1059-9495Volume 22Number 9 J. of Materi Eng and Perform (2013)22:2542-2549DOI 10.1007/s11665-013-0549-4

Crack-Growth Behavior of Laser Surface-Alloyed Low-Carbon Steel

Roman Šturm, Matjaž Žnidaršič & JanezGrum

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Crack-Growth Behavior of Laser Surface-AlloyedLow-Carbon Steel

Roman Sturm, Matjaz Znidarsic, and Janez Grum

(Submitted December 10, 2012; in revised form March 22, 2013; published online April 12, 2013)

Crack-growth behavior of Nd:YAG laser surface-alloyed as-received low-carbon steel Fe360B was evalu-ated. Thin surface layer was alloyed with silicon carbide SiC. During laser surface alloying process SiCpowder dissolved in the melted pool. The surface-alloyed layer had as-solidified structure composed mainlyof dendrites of ferrite, fine martensite needles, and retained austenite. The micro-hardness of the lasersurface-alloyed layer was about 850 HV0.1. In laser surface-alloyed layer compressive residual stresses ofaverage amount of rRS = 2100 MPa were obtained. In crack-growth tests comparison between specimensof as-received low-carbon steel Fe360B and the same steel with laser-alloyed surface was made. As the crackpropagation was perpendicular to the interface between the laser-alloyed layers and the base metal, lasersurface-alloyed specimens exhibited higher crack-growth resistance in the low stress intensity factor rangeDKth than as-received steel specimens.

Keywords DKth, compressive residual stresses, crack-growthbehavior, laser surface alloying, SiC powder

1. Introduction

Laser surface alloying (LSA) is one of the promisingprocesses for surface modification of low-carbon and low-alloyed steels. Surface alloying is an attractive processdeveloping corrosion and/or wear resistant layers, due to theinherent rapid solidification and high concentration of keyelements, combined with a good metallurgical bonding to thebase metal. Industrial applications include mainly CO2 andNd:YAG laser sources that provide appropriate power, andconsequently, energy density required for carrying out materialsprocessing.

Case hardening steels are widely used for power transmis-sions gears due to the combination of high hardness in the caseand good toughness in the core (Ref 1, 2). In the literature,Akay et al. (Ref 3) presented the effect of heat treatment on thefatigue life of low-carbon steel. The experimental resultsshowed that fatigue limit of carbon steel increased withgrowing surface hardness and surface compressive residualstresses (RS). Farfan et al. (Ref 4) have showed, the greater thecase depth the higher the fatigue strength. Liu et al. (Ref 5)have showed that fatigue properties of case hardened steelincreased after carburization. Sirin et al. (Ref 6) and Terres et al.(Ref 7–9) have studied the effect of ion nitriding of low-alloyedsteel. It has been found that the ion nitriding surface treatmentimproves the fatigue strength and increases the fatigue limitdepending on the case depth. Compressive RS on the surface

and surface hardness results in improvements in the fatigue lifebecause the thin, hard layer prevents plastic flow. The effect ofincreasing the case depth on the fatigue limit can be viewed aseffectively moving the fatigue crack initiation site further intothe core. This means that a greater applied bending stress willbe required at the surface to create a sufficiently high level ofstress at the case-core interface to initiate failure.

Thawari et al. (Ref 10) investigated LSA of medium carbonsteel with SiC powder to improve wear and corrosion resistanceof machine parts. They studied the influence of differentconditions of laser alloying by changing the laser power andscan speed in terms of microstructure, phase analysis, andmicro-hardness achieved. Nearly complete decomposition ofthe SiC leads to the formation of iron carbides and iron silicidessuch as Fe-C-Si, Fe2Si, Fe3Si, Fe7C3, and Fe3C in the remeltedlayer. Fe content in the various phases increases in the orderFe2Si fi Fe3Si fi Fe7C3 fi Fe3C, which is connected withrepetitive scanning.

Woldan et al. (Ref 11) successfully produced LSA layers oflow-carbon steel with SiC powder. Microstructural examinationof the substrate confirmed that in the laser-alloyed zone needle-like martensite is obtained, and retained austenite (RA) is foundbetween the needles of martensite. Martensitic structure variedin the needle size. The martensitic needles formed in the earlystage of austenite to martensite transformation were very large,and because of that their ends represent crack initiation places.Another similar, but yet so different way to improve wearresistance of the surface is to implement the tungsten inert gas(TIG) arc-processing on SiC surface alloying of low-carbonsteel. Buytoz (Ref 12) successfully implemented TIG processfor surface alloying, where SiC particles completely dissolvedduring the process. Also in this case, dendritic structure wasobserved in the coating layer which resulted from the rapidsolidification. Another reason for obtaining the dendriticstructure is the large difference between the melting points ofiron and other phases formed by SiC dissolution.

Grum and Znidarsic (Ref 13) focused the research of LSAmainly on highly accurate production technique and on

Roman Sturm, Matjaz Znidarsic, and Janez Grum, Faculty ofMechanical Engineering, University of Ljubljana, Askerceva 6, 1000Ljubljana, Slovenia. Contact e-mails: roman.sturm@fs.uni-lj.si,matjaz.znidarsic@fs.uni-lj.si, and janez.grum@fs.uni-lj.si.

JMEPEG (2013) 22:2542–2549 �ASM InternationalDOI: 10.1007/s11665-013-0549-4 1059-9495/$19.00

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selection of alloying powders to get suitable mechanicalproperties and RS in the surface-modified layer. In that papervery little effort has been devoted to research on mechanicalproperties such as fatigue behavior.

In LSA, laser melts a shallow region of the base metalsimultaneously with the SiC additive which ensures alloying ofthe molten base metal. McDaniels et al. (Ref 14) showed that ifthe additives are the right ones the fatigue resistance is almostunaffected by this process.

Tsay et al. (Ref 15, 16) found out, that laser surface-remeltedspecimens with no additional welding material exhibited ahigher resistance to crack growth in the low stress intensityfactor range DK than the as-received steel plates. Fractographicexaminations indicated the existence of trans-granular andinter-granular fracture modes in the surface-remelted metal.

Yang et al. (Ref 17) showed a mechanical model of acoating/laser pre-quenched steel substrate specimen. The crackwas oriented perpendicular to the interface between the coatingand the hardened layer to quantify the effects of the residualstress and hardness gradient on the crack driving force. Thiswork suggests that higher residual compressive stress, higheryield strength, and higher Young�s modulus of the hardenedlayer can greatly decrease the crack driving force. The decreasein the crack driving force means an increase in the resistance tocrack growth, and this situation can improve the fracturebehavior or improve the load-bearing capacity of the material.High residual compressive stress and high strength resultingfrom laser heat treatment also play an important role inimproving fatigue behavior, such as the decrease in the fatiguecrack-growth rate or the increase in resistance to crack growth.

Therefore, the main aim of this paper is to evaluate crack-growth behavior of the LSA steel. The prime objectives indevelopment of advanced surface layers are to improve thesurface quality and to extend the life of the component. Toachieve these objectives, it is essential to understand fractureproperties which directly affect the life of altered surfacecomponent.

In this investigation, an experimental study was conducted toevaluate LSAeffects on the crack-growth behavior of low-carbonsteel. The mechanical properties of SiC-alloyed structures wereinvestigated by using a bending fatigue machine, a micro-hardness tester, a hole-drilling residual stress measurementmethod, and optical and scanning electron microscopy (SEM).

2. Experimental Procedure

An as-received common structural steel Fe360B (ISO), i.e.,EN S235JR, was selected for LSA. Chemical composition of

the alloy (in wt.%) is presented in Table 1. The minimum yieldstrength Reh of such steel is 235 MPa, the tensile strength Rm

between 360 and 510 MPa, and hardness 180 HV3 (Ref 18).For the study, specimens with a thickness of 8 mm and

surface area of 609 60 mm2 were prepared by machining.Prior to LSA the specimens were ground to ensure a meanarithmetic roughness Ra varying between 0.12 and 0.18 lm andsurface waviness below 0.02 mm. For the experiment, alloyingmaterial SiC in powder state was chosen. Hardness of alphaSiC powder was between 2450 and 2600 HK0.1. Grain sizedistribution of SiC powder is presented in Table 2. The averagegrain size of the SiC powder was below 110 lm with achemical composition of 97% SiC, the rest was C. SiC powderwas placed on a varnish on the surface of steel specimen. Thethickness of pre-placed SiC powder was between 0.1 and0.15 mm. Laser beam was then irradiated on the specimen. Thesurface of the base material was simultaneously remelted withpre-placed SiC powder during LSA process.

The surface alloying process was made using 1 kW CWCO2 laser at laser-beam power of 600 W with a spot diameteron the surface of 2.5 mm (Q = 120 W/mm2), which ensuredcomplete dissolution of SiC powder in the melt pool. Scanspeed of laser-beam movement vb = 400 mm/min ensured awidth of the remelted trace on the specimen of 1.5 mm and awidth-to-depth ratio of the remelted trace equal to 6:1 (Ref 13).Laser-beam movement was in a zig-zag pattern where laserbeam changed its direction of movement outside the specimen.The overlapping of the remelted traces was 1/3 of the remeltedwidth. Argon was used as a protecting gas with a gas flow of5 L/min. The test specimens were LSA with SiC on twoopposite surfaces. After LSA the test specimens(609 609 8 mm3) were cut parallel to treatment to 10 mmwide specimens. In these specimens a 2 mm deep V-notch wascut perpendicular to LSA layers to get Charpy-V specimens(609 109 8 mm3).

Figure 1 shows a schematic illustration of LSA experimentand specimen preparation for the crack-growth tests.

Charpy-V specimens were subjected to bending crack-growth tests in as-received steel condition, and also in as-LSAcondition. All crack-growth tests were carried out on an electro-magnetic resonance testing machine (Cracktronic�, RumulRussenberger Prufmaschinen AG, Switzerland). The programpackage allowed computer to control the fatigue tests, and thecrack length was measured with a Fractomat� device. Mea-surements of crack propagation were made according to ASTME-647 standard. The tests were performed in the regime ofdecreasing stress intensity factor DK at constant stress ratio ofR = 0.1. The stress induced in the specimen, and the high cyclicmovement of the specimen was controlled by the Cracktronicdevice. The resonance testing frequency was about 180 Hz.

Table 1 Chemical composition of steel Fe360B in weight percentage

C Mn Si Cr Ni Cu P S Al Fe

0.06 0.22 0.16 0.15 0.27 0.31 0.011 0.004 0.007 Bal.

Table 2 SiC powder size distribution

Grain size, lm 40-60 60-80 80-100 100-120 120-140 140-160 160-180 180-200

Vol.% 36 26 12 7 5 3 8 3

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Figure 2 shows a schematic illustration of bending crack-growth test experiment. Measurements of crack length andcrack growth were done with a help of resistance strain gagesRMF-A5, which were glued over the V-notch. Fatigue crack-growth tests were performed on specimens of as-received steel,and on specimens with two opposite laser-alloyed surfaces.When a crack initiates and propagates the special resistantstrain gage RMF-A5 tears apart. The change in electricalresistance measured with Fractomat corresponded to the cracklength.

The Rumul program package enables computer to collectthe electrical potential drop change in resistant strain gage, andthe elapsed number of cycles. Additionally, the overall cracklengths were measured after the tests by optical microscopy onthe specimen surface. Crack lengths measured by opticalmicroscopy were compared with the obtained values byresistant strain gages, and were in good agreement.

3. Results and Discussion

3.1 Dimensions and Micro-Hardness Profile of the Surface-Alloyed Layer

The cross-section of one single laser-alloyed trace wassemicircular, 0.25 mm in depth and 1.5 mm in width. Toachieve fully alloyed surface in LSA process, the overlappingof the alloyed traces was 1/3 or 0.5 mm of their width.Figure 3(a) shows the cross-section of specimen after theprocess of LSA. Surface-modified layer consists of 250 lmthick alloyed layer, which is followed by heat-affected zone(HAZ), and then base material. HAZ consists of two differentzones. Upper part of HAZ is rich in carbon, which diffusesfrom surface-alloyed layer. Such carbon enrichment is notnoticed in the lower part of HAZ. Figure 3(b) shows one ofmicro-hardness variations through the modified layer after LSAof SiC. The results of micro-hardness measurements indicate,that the micro-hardness variation in laser-alloyed layer is almostconstant, i.e., between 834 and 858 HV0.1. Then, micro-hardness quickly decreases in the HAZ. The region of HAZ, inwhich carbon diffuses, has micro-hardness values between 350and 550 HV0.1. In a greater depth, in which carbon does noteffectively diffuse, micro-hardness decreases to 270 HV0.1 andfurther on to around 180 HV0.1 in base as-received material.

3.2 Microstructure of the Surface-Alloyed Layer

The microstructure of base metal was mainly ferrite. LSA ofthe as-received low-carbon steel with pre-placed SiC powdercaused complete dissolution of SiC particles in the melt pool.Due to very high cooling rates the remelted surface layersolidified under non-equilibrium conditions. Figure 3(a) showsthe cross-section of LSA layer. An x-ray diffraction (XRD)

Fig. 1 Schematic illustration of LSA experiment, and Charpy-Vnotch specimen preparation

Fig. 2 Schematic illustration of bending crack-growth testingexperiment

Fig. 3 (a) Microstructure of the surface modified layer and (b) cor-responding micro-hardness measurements across the processed thick-ness

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analysis of the LSA surface of 69 12 mm in size was made ona standard Siemens D5000 diffractometer. Prior to the analysis,the LSA surface was polished and cleaned to prevent as muchas possible the influence of undesired oxides on measurements.The x-ray tube was operated at 40 kV and 30 mA. Diffractiondata were collected over a range 30-90� 2h with Cu-Ka

radiation and step width of 0.03� 2h. The results of the XRDanalysis confirmed the presence of various new phases in theLSA layer. Figure 4 shows a diffraction diagram confirming thecomplete disintegration of SiC. It was established after aquantitative analysis from the diffraction diagram using Riet-veld�s method that in the LSA layer there was 65-80% ofmartensite (a¢), 10-20% of suessite (Fe3Si), and 5-10% of otherphases. Among the other phases, cementite (Fe3C and Fe7C3)and RA seem to be most probable, but it is characteristic of thetwo phases that they are very textured or deformed since in thediffraction diagram some peaks required are missing (Ref 19).Also other authors (Ref 10–13, 20, 21) described microstruc-ture of LSA steel with SiC powder consisted of ferrite,martensite, RA, iron carbides, and silicides. Figures 5 and 6show two different morphologies of the microstructure origi-nated after LSA. Characteristic microstructure found in themiddle of laser-alloyed layer, at depth of 100 lm, is shown inFig. 5. The occurrence of large martensitic needles is the resultof very high cooling rates in the melt pool causing thediffusionless transformation. Figure 6 shows characteristicmicrostructure at the bottom of the LSA layer, at depth of200 lm. The obvious differences in the microstructure in Fig. 5and 6 can be explained by inhomogeneous distribution of

carbon from melted SiC in the melt pool and the difference incooling rates (Ref 11, 12).

3.3 Residual Stresses

Residual stresses were calculated via strain measured withstandardized hole-drilling method in accordance with ASTME837-08. During hole-drilling, the surface stress relief wasmeasured with a typical type B three-element hole-drilling-rosette. RS were calculated from the measured strains using theintegral method (Ref 22). Figure 7 shows the measurementresults of RS at different depths. The principal RS in the LSAlayer were compressive in the range of rRSmax = �52 to�124 MPa and rRSmin = �75 to �138 MPa. At the transitionfrom LSA layer to HAZ (0.25 mm depth) the character ofprincipal RS changed from compressive to tensile. Maximumtensile RS were found in the depth of 0.4 mm, i.e.,rRSmax = 200 MPa and rRSmin = 163 MPa. The maximumtensile RS then reached a semi-stable pattern.

3.4 Crack-Growth Behavior

Figure 8 shows results of crack-growth behavior of testedspecimens obtained in fatigue bending test. Figure 8(a) shows

Fig. 4 XRD analysis of LSA surface

Fig. 5 Microstructure in the middle of the LSA layer

Fig. 6 Microstructure at the bottom of the LSA layer

Fig. 7 Variations of RS through LSA layer and HAZ

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two crack-growth curves for as-received low-carbon steelFe360B (steel) and Fig. 8(b) for the same steel altered withLSA procedure (LSA steel). In the case of LSA steel, there is0.25 mm thick laser-alloyed layer on both sides of 8 mm thicksteel specimen, and crack was growing perpendicular to theinterface between the lasers alloyed layer and the base metal.Simultaneously, crack-growth behavior was measured by straingages. The crack-growth fatigue tests started at high stressintensity factor DK. The fatigue tests operated in the regime ofdecreasing DK at constant stress ratio R = 0.1. The fatiguecrack-growth rate (da/dN) versus the variation in stressintensity factor range (DK) was then plotted in log-logdiagrams, where da/dN was evaluated from the slope of a-Ncurve.

In Fig. 8(a) results of bending crack-growth tests of Charpyspecimens of as-received steel Fe360B are shown. Thresholdstress intensity factor range DKth = 6.1-6.6 MPa m1/2 wasobtained. Curves of crack-growth behavior are similar for alltest specimens of as-received steel Fe360B differing mainly inDKth. From Fig. 8(b) it can be observed that the crack-growthcurves of LSA steel specimens are shifted to the right comparedto the curves of as-received steel Fe360B specimens. The hardsurface layer with compressive RS plays an important role toretard the crack threshold stress intensity factor range DKth.This value was higher in the case of LSA steel, which was in

the range DKth = 7.4-7.6 MPa m1/2. In Paris regime, crack-growth curves were more scattered in the case of LSA steel.The brittleness of the surface-modified layer of LSA steelinfluenced unstable crack propagation in comparison to as-received steel in Paris regime. At high stress intensity factorrange the equivalent interface fracture toughness, DKIC� 17MPa m1/2, was evaluated, which was very similar for as-received steel and for LSA steel.

Figure 9 shows additional data about the crack-growthbehavior in the regime of decreasing stress intensity factorrange DK. Two characteristic crack-growth curves are pre-sented, one for as-received steel Fe360B and one for LSA steel.In the case of as-received steel, crack initiated at the tip ofCharpy-V notch and started to grow slowly at stress intensityfactor range DKIC = 16.4 MPa m1/2. In the case of LSA steel,at DKIC = 16.8 MPa m1/2 suddenly appeared a long initialcrack. Figure 10 shows crack-growth behavior versus numberof cycles for the same test results as presented in Fig. 9. In thecase of LSA steel, a long initial crack of 0.25 mm initiated atN = 9.89 104 cycles. In the case of as-received steel Fe360B,crack initiated at N = 3.89 102 cycles. The crack length wasmeasured relative to the tip of the 2 mm depth V-notch. It wasfound out that average range of interface fracture toughnesswas DKIC = 16.4-16.7 MPa m1/2 for as-received steel andDKIC = 16.8-17.4 MPa m1/2 for LSA steel. Besides, in LSAsteel crack initiated at higher number of elapsed cycles than inas-received steel.

The influence of the cross-section microstructure on crack-growth resistance in Paris regime is described in Eq 1.

da=dN ¼ C ðDKÞm ðEq 1Þ

where da/dN (mm/cycle) is the crack-growth rate; DK(MPa m1/2) the stress intensity factor range; and C and m arethe material�s constants.

The results of such an equation can be used to predict thecrack-growth life of a component. Parameters of the crack-growth rate model for as-received steel Fe360B and for LSAsteel are presented in Table 3. DKth and DKIC were determinedfrom experimental results presented in Fig. 8 and 9, andmaterial constants C and m are calculated according to data inParis regime. The crack threshold stress intensity factor rangeDKth increased from 12 to 24% in LSA steel by comparison to

Fig. 8 Relationship between da/dN and DK for as-received steelFe360B (a) and for LSA steel (b)

Fig. 9 Relationship between stress intensity factor DK vs. numberof cycles N

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as-received steel specimens. The interface fracture toughnessDKIC increased from 0.5 to 6% in LSA steel by comparison toas-received steel specimens. In the Paris regime materialconstant m increased from 28 to 66% in LSA steel bycomparison to as-received steel specimens. LSA of SiCparticles significantly influenced crack threshold stress intensityfactor range DKth. But when a crack initiated at the surface, itpropagated faster due to material�s brittleness, as it is evidentwith the increase of coefficient m (Ref 17). The results ofbending fatigue strength by other authors (Ref 23) showed thatthe fatigue limit was associated with the microstructure, thecase depth, the distribution of RA, and the compressive RS inthe surface modified layer. Sirin et al. (Ref 6) found a linearrelationship between fatigue strength of steel and case depth.

The ratio of DKth/DKIC represents the brittleness in crack-growth behavior. In other words, this ratio in brittle metals is ofhigh magnitude, while the ratio in ductile materials is ofrelatively low magnitude. According to results in Table 3, theratio of DKth/DKIC was 0.44 in the LSA steel, and it was about15% higher than of as-received steel (0.38).

3.5 Fatigue-Fractured Surfaces

The surfaces of the fatigue-fractured specimens wereexamined by means of an optical and a scanning electronmicroscope (SEM). In Fig. 11 fatigue-cracked surfaces of as-received steel Fe360B and of LSA steel specimens are shown.Figure 11(a) presents a parabolic crack propagation front acrossthe specimen, i.e., faster crack propagation in the center than atthe surface of bending specimen. Fatigue failure in the high-cycle regimes is often dominated by the crack initiationprocesses, which are strongly influenced by the salient featuresand defects in the microstructure. Competing fatigue mecha-nisms involving crack initiation at persistent slip-bands, grain

boundaries, pores, and non-metallic inclusions or particles,have been reported to occur at surface sites. Surface finish has astrong influence on fatigue crack initiation and RS due tosurface preparation can have considerable influence on thenumber of cycles to initiation (Ref 24). At the smooth surfaceof V-notch tip we can notice small indents, steps, and gapsproduced during regular cutting process of V-notch. Thesedefects could present origins for crack initiation and propaga-tion seen as striations on the fatigue-fractured surface. Thesurface of V-notch tip of LSA specimen was very similar inquality in comparison to as-received steel. In LSA specimen,propagation front is almost perpendicular to the surface-alloyedlayer (Fig. 11b). That means initiation and propagation ofthe as-received metal in the center of specimen was at thebeginning retarded by hard LSA layers on both sides of thespecimen.

Figure 12 shows a SEM figure of a fatigue-cracked surfaceof LSA steel specimen. High hardness of LSA layer restricteddeformation of the specimen. In the surface-alloyed layer

Fig. 10 Crack-growth behavior vs. number of cycles (a-N curve)

Table 3 Parameters of the crack-growth model

Material DKth, MPa m1/2 C m DKIC, MPa m1/2

As-received steel Fe360B 6.1-6.6 (4.1-25)9 10�11 4.2-5.3 16.4-16.7LSA steel 7.4-7.6 (2-6.2)9 10�13 6.8-7 16.8-17.4

Fig. 11 Fatigue-cracked surfaces after 39 106 cycles, decreasingDK, R = 0.1. (a) As-received steel Fe360 B and (b) LSA steel

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compressive RS between 160 and 200 MPa could retard crackinitiation. A remarkable improvement in fatigue properties ofthe steel was reported due to the formation of compressive RSin the case region, which increased surface hardness (Ref 6–9).

This all influenced the rise of LSA specimen fracturetoughness stress level to DKIC = 16.8 MPa m1/2. After9.89 104 cycles at this stress level suddenly originated initialcrack at the surface of LSA layer, length of 0.25 mm, asindicated in Fig. 10. After the initiation, the crack propagatedfast through a brittle surface-alloyed layer with trans-granularcleavage and tearing mechanisms. Fast and unstable growth ofa crack originated in stress intensity level from DKIC = 16.8MPa m1/2 to DK = 14 MPa m1/2, when the crack reached0.7 mm of depth. With additional decreasing of DK, the crack-growth process became stable in Paris regime.

In Fig. 13 fracture surfaces of LSA steel are shown. Highhardness and compressive RS of the LSA layer effectivelyretard the crack initiation in the soft as-received steel at theearly stage of the crack-growth process. But when the crackinitiated in the LSA layer, the compressive RS in the LSA layercould not effectively reduce the crack-growth rate anymore. Atthat time the brittleness of the hard surface layer dominated inthe crack-growth process. This difference in crack propagationrate resulted in striations at the fractured surface in the as-

received material, where slow propagation of base metal metfaster crack propagation of LSA surface. The fracture mech-anism in the LSA layer was mainly trans-granular.

4. Conclusions

This research was done to understand LSA of SiC powderon crack-growth behavior of as-received steel Fe360B. Theconclusions reached as the result of this review are as follows:

• LSA of the low-carbon steel with pre-placed SiC powdercaused complete dissolution of SiC particles in the meltpool. The depth of cross-section of LSA layer was0.25 mm.

• Results of microstructural research showed that in theLSA layer, ferrite-martensite microstructure with someRA, iron carbides, and silicides were obtained. Micro-hardness of such a microstructure was around 850 HV0.1.

• RS were introduced into the material after uneven heatingand cooling during LSA process, which depended greatlyon remelting conditions and phase concentration. Martens-ite phase transformations in LSA layer caused an increase

Fig. 12 Fractured surface of LSA steel after fatigue crack propagation test

Fig. 13 Fatigue-fractured surface of LSA steel, SEM

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in volume and contributed to compressive RS of maxi-mum rRS = �130 MPa.

• Bending fatigue tests were performed to evaluate crack-growth behavior. Two LSA layers were on the oppositesides of the specimen, and V-notch was put perpendicularto the interface between the LSA layers and the as-received metal. The crack induced in the V-notch tip andpropagated into material perpendicular to LSA layers.Crack-growth was influenced by properties of as-receivedsteel and of thin LSA layers.

• LSA specimens exhibited higher resistance to crack initia-tion in the low stress intensity factor range DKth than as-received steel specimens. The presence of compressive RSand high hardness of LSA layers held back crack initia-tion and crack growth at the low stress intensity factorrange from values of DKth = 6.1-6.6 MPa m1/2 in as-received steel specimens to values of DKth = 7.4-7.6 MPa m1/2 in LSA steel specimens. Two 0.25 mmthick LSA layers in 8 mm thick steel specimen influencedthe increase of threshold stress intensity factor DKth in therange of 12-24%. The interface fracture toughness DKIC

increased from 0.5 to 6% in LSA steel specimens accord-ing to as-received steel specimens. In the Paris regimematerial constant m increased from 28 to 66% in LSAsteel specimens according to as-received steel specimens.

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Journal of Materials Engineering and Performance Volume 22(9) September 2013—2549

Author's personal copy