Fatigue and fracture behavior of laser clad repair of AerMet® 100 ultra-high strength steel

International Journal of Fatigue 85 (2016) 18–30

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Fatigue and fracture behavior of laser clad repair of AerMetÒ 100 ultra-high strength steel Jorge M. Lourenço a,b, Shi Da Sun b,⇑, Khan Sharp c, Vladimir Luzin d, Aloisio N. Klein e, Chun H. Wang b, Milan Brandt b a

Rio Grande do Norte Federal Institute of Technology, Industry Academic Department, Av. Sen. Salgado Filho, 1559, Tirol, Natal, RN 59015-000, Brazil Royal Melbourne Institute of Technology, Centre for Additive Manufacturing, School of Aerospace, Mechanical and Manufacturing Engineering, Carlton, VIC 3053, Australia Aerospace Division, Defence Science Technology Group (DSTG), Fishermans Bend, VIC 3207, Australia d Australian Nuclear Science Technology Organisation (ANSTO), Lucas Heights, NSW 2232, Australia e Federal University of Santa Catarina – Materials Laboratory, Mechanical Engineering Department, Florianópolis, SC 88040-900, Brazil b c

a r t i c l e

i n f o

Article history: Received 22 April 2015 Received in revised form 16 November 2015 Accepted 17 November 2015 Available online 28 November 2015 Keywords: Laser cladding AerMetÒ 100 steel Fatigue & fracture Residual stress

a b s t r a c t The effect of laser cladding on the fatigue and fracture behavior under variable amplitude loading is a major consideration for the development of laser cladding process to repair high value complex fatigue critical aerospace military components, that otherwise would be replaced. The selected material, AerMetÒ100, is a widely used ultra-high strength steel in current and next generation aerospace components, such as landing gears. Laser cladding was performed using AerMetÒ 100 powder on AerMetÒ 100 fatigue substrate specimens. No micro-cracking and very little porosity were observed in the clad layer. The fatigue tests were performed under variable amplitude loading with a maximum stress of 1000 MPa. Residual stress, microstructure, and hardness, was also evaluated. Both the as-clad and post-heat treated (PHT) samples were compared to a baseline sample with an artificial notch to simulate damaged condition. Results show that laser cladding significantly improves fatigue life, as compared to the baseline sample with a notch. However, the fatigue life of the as-clad sample is lower as compared to a baseline sample without a notch. A compressive residual stress of 300–500 MPa was observed in the clad region and HAZ. The fracture modes in the as-clad specimen consisted mainly of tearing topology surface and some regions of decohesive rupture through the columnar austenite grains. The PHT condition however was not effective in improving the fatigue life. The fracture modes showed mainly decohesive rupture, and as a consequence, reduced the fatigue life. Ó 2015 Elsevier Ltd. All rights reserved.

1. Introduction A key focus in the aerospace sector is maintaining the structural integrity of the current and next generation fleet. Ultra-high strength steels are commonly found in aircraft for example the landing gear components due to their excellent performance when exposed to high permanent and fluctuating stresses [1]. The disadvantage of ultra-high strength steels is their susceptibility to fatigue resulting in brittle fracture. The primary interest of the aerospace industry is to improve the structural integrity by repair and maintenance of parts [2–4]. Ultra-high strength steels in fatigue critical aerospace military components are very difficult to repair using conventional arc based technology and are usually replaced when the damage, such as by fatigue and wear, exceeds ⇑ Corresponding author. Tel.: +61 (0) 3 9925 4071. E-mail address: [email protected] (S.D. Sun). http://dx.doi.org/10.1016/j.ijfatigue.2015.11.021 0142-1123/Ó 2015 Elsevier Ltd. All rights reserved.

dimensional limits. Since the mid 1980’s Rolls-Royce has used laser assisted repair techniques to introduce powders alloys on gas turbine blades to optimize wear characteristics with a minimum heat affected zone and consistently good quality [5]. A similar repair process using laser cladding is explored in this paper for highstrength steels. A typical repair scenario of a damaged component requires removal of the damaged area by machining a grind-out groove, depositing of compatible material alloy using laser cladding, then finally, a post-machining process to restore its originals dimensions. The process can be used as a new route to repair fatigue critical components. Laser cladding is a process that is now offered commercially for applications in different areas such as wear and corrosion resistant coatings, part repair and additive manufacturing [6]. It uses a laser to melt and fuse a metal alloy powder to form fully dense clad layer and a metallurgical bond onto a substrate with similar properties that is chemically compatible. Advantages over other repair

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methods include, (1) controlled heat input, (2) high dimensional accuracy, and (3) great process flexibility [7–10]. These recent developments make it possible to restore damaged structures to their original mechanical and chemical properties, vastly superior to the conventional welding process used in maintenance repairs. Laser cladding is suitable to treat small areas making laser cladding an ideal candidate to repair surface defects to meet specific local service requirements. The complex fatigue behavior of the composite-like as-clad material (clad layer, clad interface, HAZ, substrate) is highly influenced by the following factors: (1) variation of microstructure and microhardness, (2) internal stress concentration sites from porosity and micro-cracking, (3) residual stresses, and (4) metallurgical bond strength at the clad interface. Sun et al. [8] in recent work evaluated the mechanical properties of laser cladding repair of AISI 4340 and AerMetÒ 100 powder, on AISI 4340 substrate in an attempt to restore and improve its fatigue properties especially addressing the airworthiness certification issue important to the aerospace industry. In their study, a certain level of degradation in the mechanical properties of AISI 4340 as-clad was found. However, by changing the material of the clad layer to AerMetÒ100, the fatigue properties were improved. Only few other studies have reported on the mechanical and fatigue behavior of laser cladding repair of primary fatigue critical steel components [7,11–14]. The tensile properties of steel clad layers have shown to produce very brittle fractures, which is attributed to the rapid cooling rates, resulting in un-tempered martensitic microstructure in the clad layer. The clad interface region also promotes the highest stress concentration, which is due to local mismatch in stiffness. The effect of laser cladding on the fatigue and fracture behavior of ultra-high strength steels has not been documented in the literature. It is the purpose of this paper to investigate the viability of the laser cladding repair process for AerMetÒ100 steel. The aims of this research study are as follow: (1) Determine if laser cladding repair can restore the geometry and fatigue properties for a representative repair scenario of AerMetÒ 100 steel. (2) Investigate the influence of the laser clad process on fatigue behavior, crack growth through the clad layer/HAZ interface, and fracture behavior when subjected to the variable amplitude. (3) Determine if a Post Heat Treatment (PHT) process, designed to restore microstructure to that optimized for high strength and toughness, can improve the fatigue properties of the repair. Microstructure, hardness, and residual stress data was also evaluated to support the fatigue and fracture analysis. The fatigue fracture surfaces were examined to establish the difference in fracture mode in the clad, heat affected zone, crack stable propagation and fast fracture region. The influence of the load ratio using fatigue marker bands recorded on the fracture surface was also assessed.

2. Experimental setup and procedure

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achieve a representative in-service hardness level of 53–55 HRC. The microstructure of the AerMetÒ 100 substrate, consists of a soft and plastically deforming ferrite (bright regions) and a hard and elastically deforming martensite (dark regions), as shown in Fig. 1a. The AerMetÒ 100 steel powder was supplied by Sandvik in the form of gas atomized particles with a particle size range of 45–75 lm, as shown in Fig. 1b (see Table 1 for composition). 2.2. Fatigue test specimens Each fatigue specimen was manufactured in accordance with ASTM standard 116 (dimensions shown in Fig. 2). The overall stress concentration factor Kt was approximately 1.04 [15]. The baseline as-damaged included a notch that represented a 0.25 mm macroscopic defect simulating a fatigue crack, (dimensions of the notch are shown in Fig. 2a). The notch was wire cut using an Electrical Discharge Machine (EDM) brand Electronica Eurocut Mark II. Three or four specimens were tested for each variable in order to obtain a statistical data confidence. Shown in Fig. 3 is the laser cladding repair process of a fatigue specimen of AerMetÒ 100. Fig. 3a shows the specimen with an artificial notch in the middle representing a macroscopic defect. Fig. 3b shows a manufactured grind-out groove (radial depth of 0.5 mm) to remove the macroscopic defect. Laser cladding was performed to re-build the geometry in and around the groove, as shown in Fig. 3c. Finally, a post-clad CNC machining was performed to remove excess clad material, as shown in Fig. 3d. For all specimens, the surface roughness measurements were: RA = 0.3 and RZ = 1.6. A total of three variables were tested: (a) AerMetÒ 100 baseline as-damaged (with a manufactured artificial notch), as shown in Fig. 2a – as a comparison to show the quality of the repair; (b) AerMetÒ 100 as-clad, as shown in Fig. 2b; and (c) AerMetÒ 100 as-clad + PHT (Post Heat Treatment), as shown in Fig. 2b, as-clad plus an additional heat treatment: standard tempering heat treatment for AerMetÒ 100 alloy at 482 °C for 5 h, in accordance with AMS standards [16]. 2.3. Laser cladding Laser cladding process was performed on a TRUMPF TruLaser Cell 7020 system. The cell is equipped with a fiber delivered 3.0 kW disk laser and a coaxial laser cladding head with a focal length of 250 mm. The coaxial powder flow had a focal distance of 8.0 mm. Helium was used as the carrier gas and argon was used as the shielding gas, at 10 and 16 L/min respectively. The optimum processing parameters, as determined from clad trials, are summarized in Table 2. These processing parameters produce a clad layer with a specific clad thickness, with sufficient dilution (mixing percentage between clad and substrate), no micro-cracking, and very little porosity (Fig. 4). AerMetÒ 100 steel is a very fracture sensitive material. Any local stress concentration will result in a premature crack initiation and reduce the fatigue life. Therefore it is crucial to carefully select the correct processing parameters in order to obtain a defect free surface, which is achieved in this research (Fig. 4).

2.1. Material details 2.4. Variable amplitude load history The AerMetÒ 100 round bars, with a diameter of 31.85 mm, were supplied in normalised overage condition from Carpenter Technology (see Table 1 for composition). The bars were then solution heat treated, by heating to 885 °C for 1 h, followed by cold treatment, 73 °C for 1 h, and finally overage tempering at 482 °C for 5 h, in accordance with AMS standard 2759/2F, to

Fatigue tests were performed under uni-axial cyclic loading with variable amplitude at room temperature. A fully-automated closed-loop servo-hydraulic mechanical test machine MTS, equipped with a 400 kN load cell was used to apply the loading history. The variable amplitude loading varied between stresses ratio

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Table 1 Chemical compositions (wt%) of the AerMetÒ 100 steel substrate and AerMetÒ 100 steel powder as provided by supplier.

Ò

AerMet 100 steel substrate AerMetÒ100 steel powder

C

Mn

Ni

S

Cr

Si

Mo

Fe

Co

0.23 0.24

0.01 0.86

11.13 11.3

0.001 0.00

3.0 3.1

0.02 0.96

1.17 1.21

Bal. Bal.

13.43 13.4

Fig. 1. (a) SEM micrograph of AerMetÒ 100 gas atomized powder and (b) optical micrograph showing the martensitic microstructure of AerMetÒ 100 substrate.

Fig. 2. Dimensions of the AerMetÒ 100 fatigue specimens (a) as-damaged baseline (b) as-clad and as-clad + PHT.

of R = 0.1 (rmin/rmax = 100 MPa/1000 MPa) and R = 0.7 (rmin/ rmax = 700 MPa/1000 MPa) under frequency of 5 Hz. Two, three and four R = 0.1 marker band pattern of 10 cycles were separated by a sequence of 500 cycle R = 0.7, as shown in Fig. 5. The total length of one block of loading was 9,590 cycles. The loading history was specifically selected to quantify the crack propagation using Quantitative Fractography (QF). The rationale for selecting the two stresses of 1000 and 700 MPa was based on the fatigue strength for maraging steels (martensiticage-hardening) such as AerMetÒ100, which ranges from 600 to 700 MPa [17,18]. Stress values ranging from 50% to 60% of the yield strength are efficient to avoid premature failure of the specimen. As a result this stress range provides a good stable crack area for QF analysis.

2.5. Fatigue crack measurement QF, using a contrasting band method, was used to measure the crack propagation in the specimens. Fatigue marker bands were inserted on the fractured surface using the specific load history shown in Fig. 5 to allow crack length measurements. The ends of the fracture surfaces from each of the fatigue specimen were cut using a Struers Discotom-100 machine with a diamond abrasive saw. The fracture surface specimens were cleaned using an ultrasonic bath for 180 s. The fracture surface appearance was imaged using a macro lens camera canon EOS 50D. The images of fracture modes were obtained using a Phillips XL30 Scanning Electron Microscope. A 30 kV electron beam and 4.0–5.0 mm spot size was used at a 10 mm working distance.

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Fig. 3. Repair process for the AerMetÒ 100 steel fatigue specimen (a) as-damaged notch, (b) CNC machine grind-out, (c) Laser cladding over the grind-out area using AerMetÒ 100 steel powder and (d) Post CNC machine to remove excess clad material.

Table 2 Laser cladding parameters. Laser power (W)

Laser spot size (mm)

Transverse speed (mm/ min)

Powder flow rate (g/ min)

Step-over width (mm)

Carrier gas flow (L/ min)

Shielding gas flow (L/ min)

800

1.3

1400

5.15

0.6

10

16

Fig. 4. Longitudinal cross-section view of AerMetÒ 100 clad layer on AerMetÒ 100 substrate.

2.6. Optical microstructure & hardness Each specimen was cross-sectioned, mounted, and polished to a 1.0 lm finish. The microstructure was revealed by etching each specimen with a 10 mL HNO3 + 20 mL HCl + 30 mL H2O solution, in accordance with standard ASTM procedure E407. Microscopic examination was conducted using a Phillips XL30 SEM. Microhardness measurements were performed using a LECO LM700AT microhardness tester. An applied load of 300 gf was held for 15 s, in accordance with standard ASTM procedure E384. 2.7. Residual stress For both the as-clad and as-clad + PHT specimens, residual stress was measured using a neutron beam diffractometer at the

Australian Nuclear Science Technology Organisation (ANSTO). A wavelength of 1.67 Å was used that produced a diffraction peak from Fe(2 1 1) reflection, resulting in a scattering angle of approximately 90°. For all measurements, each specimen was rotating continuously to eliminate possible local variations and a better representation of the specimen. The measurements were split into two parts: (1) For measurements in the clad layer and HAZ, a high spatial resolution was used with gauge volume of 0.5 mm  0.5 mm  6 mm. The measurement was taken at 0.3 mm spacing for a total of 5 in-depth points (0.7 mm HAZ/substrate + 0.5 mm clad layer). The first depth point starts 0.1 mm below the top surface of the clad, then a point every 0.3 mm through the clad layer and HAZ. Each point was

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Fig. 5. A graphical representation of the loading history applied to the specimens.

measured in sin2w-mode to reconstruct the hoop stress component under the assumption of zero radial stress. Nineteen directions in 10° steps were measured to be fitted in sin2w-analysis to obtain the hoop stress component. (2) For measurements in the remainder of the HAZ and substrate, a coarse-resolution was used with a gauge volume of 2 mm  2 mm  10 mm to measure the hoop and radial directions. The gauge volume was reduced to 2 mm  2 mm  2 mm to measure the axial direction. Eight points were measured along the diameter of the specimen in 2 mm increments. The axial stress component distribution was not measured in high resolution mode due to the limitations of the neutron experiment (necessity to work with too small gauge volume of 0.3  0.3  0.3 mm3). However, the axial stress was measured in the medium resolution of 2 mm, which is a good representative value of average stress in the clad layer and HAZ combined. 3. Results and discussion 3.1. Optical microstructure and hardness The microstructure of the AerMetÒ 100 clad layer consists of a mixture of fine columnar and equaxial grains, as shown in Fig. 6a. The formation of the clad microstructure is prior austenite grain boundary (bright colour region) and lath martensite within the boundary (dark colour region). Austenite to martensite transformation is expected due to the rapid cooling rate during the cladding process [4]. AerMetÒ 100 achieves high strength from a quench heat treatment producing Ni–Fe lath martensite structure. The secondary hardening mechanism, from tempering at 482 °C for 1–5 h, further increases its strength, while retaining its toughness, from precipitation of the M2C carbides (M = Cr, Mo and Fe) in the Ni–Fe lath martensite microstructure and dissolving cementite carbides [19]. However, the time spent at that specific temperature range during laser cladding process is too short for precipitation of the M2C carbides to occur. Therefore, only an as-quenched martensitic structure is expected to form in the clad layer with partial tempering from the re-heat caused by the overlapping clad track. The HAZ consists of coarsened lath martensite structure. A good fusion bond is also observed between the AerMetÒ 100 clad layer and substrate. The PHT showed similar grain morphology, as compared to the as-clad. Carbide formation was observed in the clad layer and HAZ (Fig. 6b), which is expected to be precipitated M2C carbides. The average clad hardness was 521 HV, which was 15% lower compared to the substrate, as shown in Fig. 7. The average clad hardness is similar to the hardness of a fully martensitic structure of AerMetÒ 100 (514 HV) [20]. An overall increase in hardness was

observed across the PHT specimen (Fig. 7). The hardness in the clad layer increased to 590 HV. The hardness increase across the PHT specimen was associated with secondary hardening reaction causing precipitation hardening of the M2C carbides within the martensite structure. 3.2. Residual stress Fig. 8a shows that a compressive residual stress between 400 and 500 MPa is produced in the hoop direction in the clad layer and HAZ of the AerMetÒ 100 as-clad specimen. The jump in stress from 450 MPa to 200 MPa near the surface of the clad layer was due to post-clad CNC machining. A transition into tension residual stress is observed in the HAZ. For both the radial and hoop stress component, a maximum tensile stress of 50–100 MPa is observed in the substrate region. For the axial stress component, an average stress in the clad layer and HAZ combined was approximately 372.5 MPa, while a maximum tensile residual stress of 150– 200 MPa is observed in the substrate region. Due to the limitations of the neutron experiment, only an average axial stress is available for the clad and HAZ combined. However, it is reasonable to assume that the axial stress distribution within the clad layer and HAZ is similar to the hoop distribution (hoop curve in Fig. 8a obtained by sin2w technique), applying ideas of the progressive deposition model for cylindrical geometry by Tsui and Clyne [21]. This model assumes that the process of deposition is symmetric is normal to the cylindrical surface. The model predicts that the hoop and axial stress are equal on the surface and the through-thickness distribution is very close within the first few millimetres. The average measured axial stress near the surface of 372.5 MPa shows that the axial component is indeed similar to the hoop stress obtained by sin2w technique (Fig. 8a). It can be expected that there is some uncertainty introduced by approximating the axial stress component with the hoop in the clad layer and HAZ, however this should not exceed experimental errors dramatically. The compressive residual stress produced in the clad/HAZ layer is due to the combination of thermal shrinkage and austenite/martensite transformation. Thermal shrinkage, constrained by the cold substrate as the deposited material cools, generates tensile residual stresses in the clad layer [22]. However, austenite to martensite transformations occurs in the clad layer, and as a result, a compressive stress is generated [23]. When the clad layer cools after it solidifies, it produces a significant amount of volumetric expansion in it which could reduce the magnitude of the tensile residual stresses produced by the thermal shrinkage, or even result in compressive residual stresses instead [23–25]. Usually, improvements in fatigue life and strength are proportionally related to the depth of the compressive residual stresses imparted by surface enhancement process [26].

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Fig. 6. SEM micrographs of the clad interface of (a) AerMetÒ 100 as-clad and (b) AerMetÒ 100 as-clad + PHT.

Fig. 7. Comparison of microhardness profile measured from the surface of the specimen.

The residual stress state remained the same after PHT (Fig. 8b), which suggests the heating temperature was too low to relieve the stresses. Post-Weld Heat Treatment (PWHT) at temperatures of 600–700 °C has been shown to be efficient in reducing the residual stresses [22,27,28]. Bendeich et al. [22] analysis of a PWHT on low pressure turbine blade showed that the heat treatment minimises the magnitude of tensile residual stresses, thereby, greatly reducing the possibility of crack initiation. Similar results are observed in [69,72]. The minimised residual stresses after PWHT are due to the tempering mechanism causing a reduction in retained austenite and plastic deformation of the reformed austenite. However, post-heat treating AerMetÒ 100 at a temperature higher than 482 °C may soften the bulk substrate and decrease its strength. 3.3. Fatigue behavior The average fatigue life of the AerMetÒ 100 as-clad improved to 184,290 cycles, as compared to the as-damaged AerMetÒ 100 baseline of 16,087 cycles (Fig. 9). The repair removed the stress concentration from the artificial notch and the clad layer was able to successfully restore the geometry and some fatigue properties. However, the fatigue life of the as-clad sample is lower, as compared to a baseline sample without a notch. The literature has experimentally shown that the fatigue life of an AerMet 100Ò baseline (without a notch) survives at least 106 cycles at a maximum load of 1000 MPa [29,30]. Both papers used R = 0.1 and uni-axial fatigue test at a range of stress levels to establish a general fatigue

life curve for AerMetÒ 100. It is noted that constant amplitude loading is used in the literature, as compared to the variable amplitude used in this paper and some differences in fatigue life may occur. The lower fatigue life of the as-clad specimen, as compared to that of the baseline without a notch, is due to the difference in microstructure formation in the clad layer and the substrate. The composition is similar for both the powder and substrate material. However, a significantly different microstructure was observed in the clad layer, where it appears to be dominated by c-Fe columanar grains (Fig. 6a). Due to the fracture sensitivity of AerMetÒ 100, any difference in microstructure formation from the original material (Fig. 1b) will degrade the fatigue properties. Ultimately, the as-clad condition is still weaker as compared to the baseline without a notch. The average fatigue life was reduced for the AerMetÒ 100 asclad + PHT specimens to 95,665, as compared to the AerMetÒ 100 as-clad specimens of 184,290 (Fig. 9). It was anticipated that differences in the residual stress state will be responsible for any difference in fatigue life. However, it was evident that the residual stress remained in the same state after PHT (Fig. 8b). The reduced fatigue life was due to the change in microstructure and increase in hardness from the PHT causing a very brittle fracture. Fig. 10 shows that marker bands were clearly identified on the fracture surface of as-damaged AerMetÒ 100 baseline specimen. The first visible marker band was clearly measured when the crack had a length of 0.1 mm. This was relative from the bottom of the

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Fig. 8. Residual stress profile for AerMetÒ 100 (a) as-clad and (b) as-clad + PHT.

Fig. 9. Comparison of fatigue life of AerMetÒ 100 specimens.

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Fig. 10. Fatigue crack surface of as-damaged AerMetÒ 100 baseline specimen showing its fatigue crack propagation (left); and a high magnification of the outlined white square (right).

Fig. 11. Crack growth data from as-damaged AerMetÒ 100 baseline, as-clad and as-clad + PHT specimens.

artificial notch. Marker bands were not clearly observed in the clad layer for both as-clad and as-clad + PHT specimens. This is attributed to the fine grained martensite structure in the clad layer. The strain hardening of a fine grain martensite in clad layer is greater than in a coarse grain martensite polycrystalline, in HAZ and substrate. This effect is detrimental to plastic deformation, and as a consequence, the fracture plane change was prejudiced in the clad layer. As a result, fatigue marker bands in it are absent with little deformation becoming invisible to characterize. The measurements of the crack length for the as-clad specimens started with fatigue life of 134,400, 164,210, and 164,210 cycles, and 185,940 cycles, for crack lengths of 0.694 mm, 0.609 mm, 0.559 mm and 0.62 mm respectively (see Fig. 11). All measurements were relative from the point of crack initiation on the specimen’s surface (for instance, see Fig. 13a). These crack lengths indicate data with very high level of confidence. In the substrate region of the as-clad and PHT specimens, the crack lengths were similar to the as-damaged baseline (Fig. 11). According to the Paris law [18], region I is the near threshold where there is no observable crack growth, characterized as non-

propagating cracks. The cracks growth measurements in as-clad samples were visible only after the 20th block of loading. This number of blocks belongs to region II in Paris law curve. The behavior in region III for both samples, as-baseline and as-clad, were the same and the fatigue crack growth rates are very high as they approach instability. This examination revealed that there were many subtle markings present in the clad layer. However, the detail shown did not clearly reveal the marker bands in this region, and hence was not included in the crack growth measurements. For all specimens, the crack initiated from the top surface of the clad layer. The cracks progressed in semi-elliptical fashion in the substrate region until a full width crack that persisted up to complete failure with a final crack length of approximately 4.5 mm.

3.4. Fracture behavior For all variables, the fracture surface consists of three different fatigue characteristic zones: stable crack growth, shear lips and fast fracture (Figs. 12a, 13a, and 14a). In addition, the macroscopic

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Fig. 12. Macrograph and SEM micrographs of as-damaged AerMetÒ 100 baseline after failing at 16,087 cycles under variable amplitude at 1000 MPa maximum load: (a) fracture surface appearance, (b) location of microscopic crack initiation, (c) fracture mode in stable crack propagation, (d) fracture mode in fast fracture zone.

morphology is at 90° to the main stress axis with a flat appearance, which is consistent with a brittle fracture. For as-damaged AerMetÒ 100 baseline specimen, the significant reduction in the fatigue life was due to (1) the manufactured notch causing a significant increase in stress concentration factor, and (2) its fracture sensitive nature. As a result, multiple initiation of the main crack occurred, as shown in Fig. 12b, which was the main source of crack initiation. The stable crack growth region was flat, primarily transgranular, with characteristic fatigue marker bands due to the applied loading history. High magnification observation of the stable crack region, shown in Fig. 12c, revealed a mixture of Tearing Topology Surface (TTS) and poor shallow fatigue striations (indicated by number 1). According to Manigandan [30], fatigue striation features are similar to localized ductile microsplastic and brittle deformation, and are usually observed in the stable crack region. However, according to Wulpi [31], fatigue striations are usually not present on very hard or very soft metals. Hardened steels, such as the AerMetÒ 100, develop either poorly formed or no striations due to their lack of ductility. As hardness increases, striations are more difficult to observe, often resulting in a fracture surface that appears to be featureless, which is evident in Fig. 12c. The TTS fracture appears to be formed by ductile or microplastic tearing on a very fine scale, even, in materials with low ductility such as AerMetÒ 100. Similar fracture modes of TTS were observed in a martensitic HY130 high strength steel [32]. Overload failure by Microvoid Coalescence (MVC) was observed in the fast fracture zone with cuplike depressions, as shown in Fig. 12d. The microvoids nucleate from regions of localized strain discontinuity, associated with homogeneous distribution of nanoscale coherent M2C alloy carbide in a martensite matrix [33,34]. MVC is indicative of localized ductile failure mechanism. However,

for ultra-high strength steel, due to their very low plasticity, the MVC observed is associated with cleavage fracture mode regardless of the transgranular fracture path. Shown in Fig. 13a, the clad layer and HAZ are visible in the stable crack growth region for the as-clad AerMetÒ 100. The crack initiated from a micro defect approximately 20 lm from the top surface of the clad layer, as shown in Fig. 13b. The fracture modes in the AerMetÒ 100 clad layer were a mixture of decohesive rupture (number 2 in Fig. 13c) and TTS (number 3 in Fig. 13c). Only TTS was observed in the HAZ (number 4 in Fig. 13d) and in the remainder of the stable crack propagation (Fig. 13e), similar to the fracture mode in the as-damaged baseline specimen. MVC with very little plasticity was observed in the fast fracture zone, also similar to the as-damaged baseline (Fig. 13f). Decohesive rupture in the clad layer is believed to be the result of the weak nature of columnar grain growth and grain boundary micro-segregation by chemical impurities in the AerMetÒ 100 steel powder (Mn, Si, and S) [31,35,36]. Decohesion along columnar grain boundaries is typically observed in other cast structures such as those produced in the conventional weld process [37]. For AerMetÒ 100 as-clad + PHT specimens the crack initiation, outlined by a white square in Fig. 14a, and highly magnified in Fig. 14b, shows that crack initiated in a micro defect approximately 20 lm from the top surface of the clad layer. The same pattern was observed in as-clad specimens. Fig. 14c revealed that the clad layer exhibited mix modes of fracture: mainly decohesive rupture (number 5) and some TTS (number 6). The low fatigue life of the PHT specimens was associated with their brittleness, which is evident by the increase in hardness (Fig. 7). An increase in decohesive ruptures was also observed, as compared to the as-clad specimens. Fig. 14d shows the clad interface where the crack front transitions

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Fig. 13. Macrograph and SEM micrographs of as-clad AerMetÒ 100 after failing at 184,290 cycles under variable amplitude at 1000 MPa maximum load: (a) fracture surface appearance, (b) location of microscopic crack initiation, (c) fracture modes in clad layer, (d) HAZ layer, (e) stable crack propagation, and (f) fast fracture region.

from decohesive rupture through the columnar austenite grains (number 5 in Fig. 14d) to TTS in the HAZ (number 6 in Fig. 14d). The fracture mode in the stable crack growth region was TTS (number 6 in Fig. 14e) with evidence of poor fatigue striation (number 7 in Fig. 14e). TTS is very complex microvoids coalescence where strain localization prevents the observation of well-developed voids [32]. Only some microvoids can be noticed on TTS fracture modes (Figs. 13e and 14e). In the fast fracture zone for the asclad + PHT specimens, the main fracture mode mechanism was also microvoids coalescence with very little plasticity (Fig. 14f). All three variables tested showed similar fracture behavior in these regions. Crack growth models of fatigue assume that the crack path behavior is a combination of cleavage and plastic blunting [38]. The micrograph that represents the stable crack region, in Fig. 15a, shows two kinds of fracture mode: one with TTS tearing

(number 8) and cleavage behavior (number 9). However, the micrograph that represents the fast fracture region, see Fig. 15b, shows only cleavage fracture behavior. The stable crack region in AerMetÒ 100 steel is controlled by the cleavage fracture mode. Nanoparticles-nucleated dimples of coherent M2C alloy carbide in a martensite matrix of AerMetÒ 100 are indicated by number 10 in a microplastic area shown in the bottom left corner detail of the Fig. 15a. These microplastics regions can be associated with TTS. As discussed earlier, the MVC observed in the fast fracture region is associated with cleavage fracture mode (Fig. 15b). Cleavage is a mode of separation that occurs very suddenly between one face of the cell and the mating face of adjacent one, and it is caused by overload fracture [31,35,36]. Due to imperfections and changes in crystal lattice orientation, this mode of failure produced distinct

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Fig. 14. Macrograph and SEM micrographs of as-clad + PHT AerMetÒ 100 after failing at 95,665 cycles under variable amplitude at 1000 MPa maximum load: (a) fracture surface appearance, radially outward, (b) location of microscopic crack initiation, (c) fracture modes in clad layer, (d) in HAZ layer, (e) in stable crack propagation, and (f) in fast fracture.

cleavage fracture features in the AerMetÒ 100, such as: cleavage steps (number 11), grain or subgrain boundary (number 12), river pattern (number 13), and dimples with low plastic deformation (number 14). 3.5. Influence of stress ratio, R = 0.1 and R = 0.7 Fig. 16 shows SEM micrographs of the fatigue fracture surface where the crack plane change is caused by groups of high stress ratio (R = 0.7) and low stress ratio (R = 0.1) loading cycles. These features are a result of the slip band formation when associated with contrast between two different stress ratios produced in variable amplitude sequences. The contrast between different stress ratios is less visible during SEM examination, particularly, when the crack is small. However the bands are still there and use of

optical contrast is the best way to characterize them [39]. The crack path, texture and direction change when the crack is generated with high number of cycles at R = 0.1 interspersed with few number of cycles also at R = 0.1, but separated by high number at R = 0.7, and then the crack changes back again to the original texture but on a different crack path, as shown in Fig. 16a. White at al [38] observed these crack path changes in aluminum under periodic under loads, with R = 0.1, and phenomena such as striations, ridges, fissures ledge faces and crack forking were registered. The general macroscopic fracture surface appearance even though with the constant changing of the crack angle is normal to the main stress loading direction, however, the local sections of the crack are tilted either up or down with respect to this direction [38,40]. It has been observed in AerMetÒ 100 fracture mode that the martensitic structure failed by creation of slip band with a little

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Fig. 15. (a) SEM micrographs in a very high magnification showing an area of the stable crack region, (b) fast fracture region for AerMetÒ 100 as-clad specimen.

Fig. 16. (a) SEM micrographs of AerMetÒ 100 as-clad fatigue fracture caused by high stress ratio groups, (b) high magnification view, (c) high magnification of R = 0.1 region, (d) high magnification of R = 0.7 region.

plasticity for both stress ratios, R = 0.1 and R = 0.7, characterized by the change in crack path. These features illustrated in Fig. 16a and b indicate that single tension load with different stress ratios produces crack growth in two distinct directions and, as a consequence, slip band along two different crystals planes. So, if the crack is growing on the main plane [1 1 1] then it is generally with a slight slope either heading down, or uphill, in relation to this main plane, as shown in Fig. 16a. Fracture modes for both stress ratios, showed in Fig. 16c and d, suggest that the same fracture behavior happened in these specimens, however, with different texture orientation and these features combined with different crystals planes produce the marker bands seen in SEM images in Fig. 16a. Surfaces produced by R = 0.1 are brightly lit whereas those ones produced by R = 0.7 are all shadowed areas suggesting that one plane is reflective and another one is non-reflective or reflecting in two different directions.

4. Conclusion In this study, laser cladding of AerMetÒ 100 powder on AerMetÒ 100 substrate was performed to demonstrate viability of repair. A variable amplitude fatigue load was used for 1000 MPa maximum load and 0.1/0.7 load ratios. The fatigue and fracture behavior is studied. The following conclusions are drawn from this study: (1) Laser cladding successfully re-built the geometry in a grindout region with a defect free clad layer and a small HAZ. The average fatigue life of the as-clad specimen was improved significantly to 184,290 cycles, as compared to the as-damaged baseline condition of 16,087 cycles. However, the fatigue life of the as-clad is lower, as compared to a baseline without a notch. The lower fatigue life is due to the difference in microstructure produced in the clad layer and the substrate.

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(2) The fracture initiated from the top surface of the clad layer. The fracture modes consisted mainly of tearing topology surface with areas of decohesive rupture through columnar grains. (3) The average fatigue life of the as-clad + PHT specimen improved only to 95,665 cycles, as compared to the asdamaged baseline condition of 16,087 cycles. The low fatigue life for the PHT specimens was associated with their brittleness, which is evident by the increase in hardness and decohesive ruptures observed in the clad layer. (4) In regards to the influence of stress ratio, R = 0.1 and R = 0.7, the crack path, texture and direction change when the fatigue test is generated using different stress ratio so that the fatigue marker bands become visible by contrast. The fracture modes do not change at the different stress ratio. Acknowledgements The authors would like to acknowledge the financial support from the Brazilian Research Funding Support – CAPES/‘‘Ciências Sem fronteiras” program – process number registration BEX 9296/13-1. The authors would also like to acknowledge the technical support from the RMIT Microscopy and Microanalysis Facility (RMMF) and to Alan Jones for his assistance with laser cladding. The authors wish to record their sincere thanks to Dr Simon Barter from DSTG and Sebastian Naselli from RMIT for their important contributions. References [1] Campbell FC. Manufacturing technology for aerospace structural materials. 1st ed. Amsterdan: Elsevier; 2006. [2] Fallah V, Alimardani M, Corbin SF, Khajepour A. Impact of localized surface preheating on the microstructure and crack formation in laser direct deposition of Stellite 1 on AISI 4340 steel. Appl Surf Sci 2010;257:1716–23. [3] Alimardani M, Fallah V, Khajepour A, Toyserkani E. The effect of localized dynamic surface preheating in laser cladding of Stellite 1. Surf Coat Technol 2010;204:3911–9. [4] Bhattacharya S, Dinda GP, Dasgupta AK, Mazumder J. Microstructural evolution of AISI 4340 steel during direct metal deposition process. Mater Sci Eng, A 2011;528:2309–18. [5] Draper CW, Mazzoldi P. Laser surface treatment of metals. 1st ed. Dordrecht: Martinus Nijhoff Publishers; 1986. [6] Vilar R. Laser cladding. J. Laser Appl. 1999;11:64–79. [7] Ganesh P, Moitra A, Tiwari P, Sathyanarayanan S, Kumar H, Rai SK, et al. Fracture behavior of laser-clad joint of Stellite 21 on AISI 316L stainless steel. Mater Sci Eng, A 2010;527:3748–56. [8] Sun SD, Liu Q, Brandt M, Luzin V, Cottam R, Janardhana M, et al. Effect of laser clad repair on the fatigue behaviour of ultra-high strength AISI 4340 steel. Mater Sci Eng, A 2014;606:46–57. [9] Alam MM, Kaplan AFH, Tuominen J, Vuoristo P, Miettinen J, Poutala J, et al. Analysis of the stress raising action of flaws in laser clad deposits. Mater. Des. 2013;46:328–37. [10] Hazra M, Mondal AK, Kumar S, Blawert C, Dahotre NB. Laser surface cladding of MRI 153M magnesium alloy with (Al+Al2O3). Surf Coat Technol 2009;203:2292–9. [11] Imran MK, Masood SH, Brandt M, Bhattacharya S, Mazumder J. Direct metal deposition (DMD) of H13 tool steel on copper alloy substrate: evaluation of mechanical properties. Mater Sci Eng, A 2011;528:3342–9. [12] Niederhauser S, Karlsson B. Mechanical properties of laser cladded steel. Mater Sci Technol 2003;19:1611–6.

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