Effect of laser clad repair on the fatigue behaviour of ultra-high strength AISI 4340 steel

Effect of laser clad repair on the fatigue behaviour of ultra-high strength AISI 4340 steel

Materials Science & Engineering A 606 (2014) 46–57 Contents lists available at ScienceDirect Materials Science & Engineering A journal homepage: www...

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Materials Science & Engineering A 606 (2014) 46–57

Contents lists available at ScienceDirect

Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

Effect of laser clad repair on the fatigue behaviour of ultra-high strength AISI 4340 steel Shi Da Sun a,c,n, Qianchu Liu b,c, Milan Brandt a,c, Vladimir Luzin d, Ryan Cottam c,e, Madabhushi Janardhana f, Graham Clark a a

School of Aerospace, Mechanical & Manufacturing Engineering, RMIT University, Bundoora, VIC 3083, Australia Aerospace Division, Defence Science & Technology Organisation (DSTO), Fishermans Bend, VIC 3207, Australia c Defence Materials Technology Centre (DMTC), Hawthorn, VIC 3122, Australia d Australian Nuclear Science Technology Organisation (ANSTO), Lucas Heights, NSW 2232, Australia e Industrial Research Institute Swinburne (IRIS), Swinburne University of Technology, Hawthorn, VIC 3122, Australia f Aircraft Structural Integrity Section, Directorate General Technical Airworthiness (DGTA), RAAF, Laverton, VIC 3027, Australia b

art ic l e i nf o

a b s t r a c t

Article history: Received 10 December 2013 Received in revised form 20 March 2014 Accepted 21 March 2014 Available online 28 March 2014

The fatigue behaviour of an ultra-high strength steel ( 41800 MPa) was evaluated to assess the potential of using laser cladding as a repair tool for such steels in aeronautical structural applications. AISI 4340 and AerMet 100 steel powder were used to clad over a grind-out region in an AISI 4340 steel substrate using a 2.5 kW ND:YAG laser. Post-clad heat treatment (PCHT) was also investigated. Results showed very poor tensile properties and significantly reduced fatigue life of the AISI 4340 as-clad with a very high hardness and brittle fracture in the clad and HAZ zone. Residual stress results showed a compressive residual stress in the clad region and tensile residual stress in the HAZ. Changing the alloy of the clad layer to AerMet 100 steel, as well as applying a PCHT process, showed promising results as the fatigue life was improved from that of the grind-out substrate. & 2014 Elsevier B.V. All rights reserved.

Keywords: AISI 4340 steel AerMet 100 steel Laser cladding Residual stress Fatigue behaviour

1. Introduction Ultra-high strength steels such as AISI 4340 and AerMet 100 are materials used widely in modern aircraft structure, particularly in critical applications such as undercarriages. However, they are highly sensitive to damage caused by corrosion, fatigue, stress corrosion cracking, and in some of these applications, to impact damage from foreign objects. High strength steels achieve their strength at the expense of toughness, and as a result, any damage which promotes crack development and propagation increases the risk of unpredicted catastrophic failure. Discovering even the smallest cracks (in the sub-millimetre scale) will usually require some attention, since the components are usually managed on a safe-life (no detectable cracks) basis. Grind-out is a technique traditionally used to remove the damage. However, aircraft components are geometrically optimised, leaving very little material to safely grind-out. This method is viable only if the grind-out does not exceed the dimensional limits. Exceeding such limits would n Corresponding author at :School of Aerospace, Mechanical and Manufacturing Engineering, Advance Manufacturing Precinct, RMIT University, Bldg 55, Level 4, 58 Cardigan Street, Carlton, VIC 3053, Australia. Tel.: þ 61 3 9925 4071. E-mail address: [email protected] (S.D. Sun).

http://dx.doi.org/10.1016/j.msea.2014.03.077 0921-5093/& 2014 Elsevier B.V. All rights reserved.

lead to replacement of the component, resulting in high cost and reduced availability. One attractive method of repair is to use laser cladding to rebuild material in a grind-out, which has exceeded the permissible limit, to restore the mechanical properties, especially fatigue life, back to a minimum safety level. Laser cladding technology uses a laser to melt and fuse a metal alloy powder, with similar or enhanced properties to the application, onto a substrate [1,2]. Other possible repair methods include cold spray, traditional arc welding, and plasma spray [3]. However laser cladding has received much more focus as a repair method due to low heat input, strong fusion bond between the added metal and the base material, and its ability to produce very little distortion and micro-cracking [2]. Traditional repair methods such as argon-arc welding usually produce deep heat-affected zones (HAZ), visible deformation, and are highly susceptible to hydrogen cracking [1]. Another major advantage of laser cladding is that the powder alloy can be optimised by changing the alloying composition of the powder to optimise the mechanical properties. Laser cladding is particularly suitable for aerospace components because they deal with larger, complex shaped components, and will offer major cost savings if the process is successful. However, repairing aero-grade high strength steel is particularly difficult since it deals with high structural loadings, fatigue critical components and a

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Table 1 Chemical compositions (wt%) of the AISI 4340 base material and AISI 4340 and erMet 100 powders, as provided by supplier.

AISI 4340 steel substrate AISI 4340 steel powder AerMet 100 steel powder

C

Mn

Ni

S

Cr

Si

Mo

V

Fe

Co

0.41 0.4 0.25

0.7 0.7 0.63

1.74 1.69 11.2

o0.02 o0.02 –

0.77 0.72 3.1

0.24 0.2 0.64

0.25 0.28 1.28

0.046 0.05 –

Bal. Bal. Bal.

– – 14.0

Fig. 1. Schematic representation of the experimental procedure for the mechanical testing of AISI 4340 steel: (a) 0.7 mm grind-out depth along the direction of rolling, (b) multi-track cladding using the optimum processing parameters to fill the grind-out area, (c) excess clad layer removed by a CNC machine for a flat surface finish and (d) individual dog-bone specimens machined by wire-cutting.

need for damage tolerance. A repair by rebuilding material to restore strength which would allow parts to remain in service, and thus may reduce operating costs, will only be possible if it does not increase the risk of cracking and degraded mechanical properties. This research evaluates the mechanical properties of laser cladding repair of ultra-high strength, in an attempt to restore and possibly increase the fatigue properties above the un-repaired case. The mechanical properties of AISI 4340 steel have been well documented [4–6] and a number of publications have reported on the mechanical properties of surface melting of AISI 4340 steel [7,8], where it has been demonstrated that the so-called HAZ does not have an adverse effect on the fatigue properties of AISI 4340 steel [8]. For laser cladding repair applications, most literature has focused on the repair of secondary components [9,10], where geometry restoration and restoring surface properties, such as wear and erosion properties, are the major concern. Few publications have reported on the mechanical properties of the repair of primary fatigue critical components [11,12], and microstructure evolution on laser cladding of ultra-high strength AISI 4340 steel [13,14]. However, none have been published on the mechanical properties and fatigue behaviour of laser cladding repair in ultrahigh strength steel aeronautical structural applications, where the substrate strength is greater than 1800 MPa. This paper focuses on the mechanical properties of laser cladding using quench and temper aero-grade AISI 4340 steel substrate. Two different alloy powders are used and a post-heat treatment process is investigated. Microstructure and microhardness evolution, static properties, fatigue S–N curve, fracture behaviour, and residual stress properties were analysed and reported. This study provides essential results on the overall potential of laser cladding as a future repair solution for aircraft landing gear.

2. Experimental setup 2.1. Material preparation AISI 4340 steel plates of dimension 200 mm  160 mm  7 mm (composition shown in Table 1) were supplied in annealed condition. The plates were then hardened by heating to 850 1C for 1 h, followed by oil quench, and finally tempered at 220 1C for 4 h, in accordance with AMS standard 2759/2 F [15], to achieve a representative in-service hardness level of 53–55 HRC. The microstructure is fine tempered martensite. AISI 4340 and AerMet 100 powders were supplied by Sandvik Osprey Ltd. in the form of gas atomised spherical particles (composition shown in Table 1), with a particle diameter in the range of 45–106 mm. A representative in-service grind-out was applied on each plate using a CNC mill. The grind-out depth was 10% of the thickness (0.7 mm 70.05 mm) with a width of 10.0 mm (Fig. 1a). A variation in grind-out depth is expected due to distortion of the plate. 2.2. Laser cladding A fibre-coupled Rofin Sinar CW025 Nd:YAG laser and a sideinjecting powder delivery nozzle were used for the cladding process. The laser optics comprised a 200 mm collimation lens and a 200 focusing lens located 204 mm from the substrate surface, producing a laser spot of about 2.5 mm. The laser spot of 2.5 mm was measured by a beam analyser corresponding to the diameter of the Gaussian profile. The powder delivery nozzle diameter was 2 mm and it was located 11 mm from the substrate surface at an angle of 251 from the normal. Argon was used as both the carrier and shielding gas at a flow of 15 L/min. The laser processing parameters used in this experiment are shown in

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Table 2. The laser clad heat input was approximately 75 J/mm. The approximate track length was 19 mm and a total of 150 overlapping clad tracks were required to fill the grind-out area (Fig. 1b). The overlapping track width was 1.0 mm. The top surface of the clad was removed by an Elliott 921 magnetic grinding machine equipped with an aluminium oxide abrasives wheel to ensure the test section of the plate was completely flat (Fig. 1c). Finally, each plate was wire-cut into four dog-bone specimens for tensile and fatigue testing (Fig. 1d). For a consistent surface finish across the specimen, the surface of the test section of all specimens was machined by removing approximately 5–10 mm off both sides of the specimen using an Elliott 921 magnetic grinding machine with a fine grade of 125–220 mm grit size. The approximate surface roughness values (measured in the longitudinal direction across the test section) were RA ¼0.2 mm and RZ ¼2.0 mm. In this study, a total of four variables were tested, as shown in Fig. 2. A traditional post-weld heat treatment (PWHT), also known as stress relieving heat treatment, was not used since the tempering would also significantly reduce the high substrate hardness and strength of the substrate. The post-clad heat treatment (PCHT) process of the clad plates was used in this study. The PCHT process was performed in accordance with AMS standard 2759/2F [15] of annealing, hardening and tempering: annealing at 830 1C for 1 h followed by slow cooling, then hardening by heating to 850 1C for 1 h, followed by oil quenching, and finally tempering at 220 1C for 4 h. 2.3. Experimental procedures The clad specimens were cross-sectioned using a diamond saw and then hot mounted in epoxy. A 250, 400, 600, 800, and 1200 grit size was used to grind each sample. An applied force of 150 N and 10–30 s was used on each grinding stage. A 9, 3, 1, and 0.25 diamond paste was used to polish each sample. An applied force of

Table 2 Laser processing parameters used for the laser cladding experiment. Laser power (kW)

Powder flow rate (g/min)

Laser traverse speed (mm/min)

Laser spot size (mm)

Overlap width (mm)

1.25

12.5

1000

2.5

1.0

150 N and 15–20 min was used on each polishing stage. Finally, a 2% Nital solution was used to reveal the microstructure of AISI 4340 and 10 mL HNO3 þ20 mL HCl þ30 mL H20 solution was used to reveal the microstructure of AerMet 100, in accordance with standard ASTM procedure E407. Microscopic examination was conducted using a Leica MEF3 optical microscope. Microhardness measurements were performed using a LECO LM700AT micro-hardness tester. For each indent, an applied load of 300 gf was held for 15 s, in accordance with standard ASTM standard E384. For comparison, the hardness of the substrate (AISI 4340) was also measured. Composition analysis was performed using a XL30 Scanning Electron Microscope equipped with EDAXs (Energy-dispersive X-ray spectroscopy). Cladding of residual stress samples was performed using the same processing parameters as the ones used for the mechanical test plates, to achieve clad dimensions of 30 mm wide and 30 mm long. Residual stress profiles were measured using a residual stress diffractometer. A wavelength of 1.67 Å and a diffraction from Fe {211} reflection were used for the monochromatic neutron beam resulting in a scattering angle of (2θ) approximately 901. A gauge volume of 0.5  0.5  20 mm3 was used for the neutron strain measurements. Fourteen depth points were measured for each of the four variables in the vertical direction through the clad, heat affected region, and substrate at a spacing of 0.5 mm for a total distance of 7.0 mm. Each of the fourteen depth points were measured in three principal directions (two in-plane directions and one normal to the surface). Axial tension testing was performed using a 250 kN MTS testing machine using a strain rate of 1.0 mm/min, in accordance with ASTM standard 8M (dimensions shown in Fig. 3a). Strain was measured using a 25 mm MTS extensometer. Four specimens were tested for each variable. For comparison, the tensile properties of both the substrate (AISI 4340) and substrate with a grind-out were also tested. Axial constant amplitude fatigue testing was performed using a 100 kN MTS testing machine using an R ratio of 0.1 and test frequency of 10 Hz, in accordance with ASTM standard E466 (dimensions shown in Fig. 3b). Testing was performed at room temperature. For each variable, three to four specimens were tested over four different stress levels to produce a preliminary fatigue life curve, in accordance with ASTM standard E739. For comparison, the fatigue life of both the substrate and substrate with a grind-out were also tested. Fracture surface images were obtained using a Phillips XL30 Scanning Electron Microscope.

Fig. 2. Schematic representation of the cross-section view of the fatigue dog-bone specimens and showing the four variables used in this study: (a) AISI 4340 as-clad, (b) AerMet 100 as-clad, (c) AISI 4340 as-clad þPCHT and (d) AerMet 100 as-clad þPCHT.

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Fig. 3. Top view and dimensions of the dog-bone specimens where the shaded area represents the clad: (a) tensile specimen and (b) fatigue specimen.

Fig. 4. Micrographs of the clad layers: (a) AISI 4340 as-clad etched with 2% Nital solution, (b) AerMet 100 as-clad etched with HNO3–HCl–H20 solution, (c) AISI 4340 ascladþ PCHT etched with 2% Nital solution and (d) AerMet 100 as-clad þPCHT etched with HNO3–HCl–H20 solution.

3. Results 3.1. Microstructure Fig. 4a and b shows that the AISI 4340 and AerMet 100 clad layer consists of fine columnar and cellular grains where the

growth is in the direction of solidification. The formation of clad microstructure is a typical columnar prior austenite grain structure due to solid state transformation, similar to steel cast structure, and is well documented in welding applications [16]. The appearance of the columnar and cellular grains varies across the clad layer caused by the varying temperature gradients that generally

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Fig. 5. Micrographs of AISI 4340 HAZ etched with 2% Nital solution: (a) clad interface, (b) HAZ 0.1 mm below the clad interface, (c) HAZ 0.6 mm below the clad interface and (d) HAZ interface.

exist across the solidifying structure resulting in the formation of different sized columnar and cellular grains [16]. Fig. 4c and d shows the microstructure of AISI 4340 and AerMet 100 clad layer after PCHT (850 1C þoil quench and 220 1C temper). The microstructure in AISI 4340 as-clad after PCHT was uniform across the specimen and consisted of fine tempered lath martensite (Fig. 4c). The microstructure in AerMet 100 as-clad after PCHT consisted of a coarsened columnar and cellular grain structure in the clad region (Fig. 4d). Fig. 5a shows the clad interface region where the microstructure changes from solidification structure to AISI 4340 HAZ. The clad interface region shows a good metallurgical bond. Fig. 5b shows the HAZ adjacent to the clad interface where the microstructure primarily consisted of coarsened lath martensite. The microstructure becomes finer further away from the clad interface, as shown in Fig. 5c. The coarsening was due to the high austenitising temperature adjacent to the clad interface, causing micro-alloy precipitates to dissolve, and unpinning of austenite grain boundaries is expected with substantial growth of prior austenite grains [17]. In Fig. 5d, the white etched region represents the HAZ where partial austenitising occurs, while the darker etched region represents the over-tempered zone of the substrate. The HAZ interface is a good indication when the heating temperature drops below the critical Ac1 temperature (727 1C). 3.2. Microhardness Fig. 6 shows the comparison of vertical microhardness profile of a cross-sectioned sample for all variables tested. In Fig. 4a, for the AISI 4340 as-clad, the region A–B, clad layer, shows that the hardness of the clad was approximately 650 Hv which was 30–40% higher than the substrate material (580 Hv). The high hardness in the clad layer is attributed to the hard and brittle properties of martensite formed at rapid cooling rates. The region B–C, HAZ, shows the hardness sharply increased from 600 Hv to 720 Hv over a distance of 1.5 mm. The increased hardness is due to the rapid

Fig. 6. Vertical microhardness profile, measured relative to the clad interface: (a) as-clad condition and (b) PCHT condition.

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self-quenching from austenite to martensite caused by the heat sink into the substrate material [7,8,18]. A similar microhardness trend in the HAZ was observed in laser melting applications of AISI 4340 steel [8]. At point C, the interface between the HAZ and the substrate, the hardness decreased sharply from 720 Hv to 380 Hv, which was a result of heating temperature not achieving the critical Ac1 temperature, indicating no austenitising had occurred. The region C–D, the part of the substrate, is known as the overtempered zone in welding applications. Overall, this region was soft and its hardness relatively lower compared to the original substrate. The cause of the softening is due to tempering, where the carbon atoms diffuse from martensite in the substrate, to form a carbide precipitate resulting in the formation of ferrite and coarsened cementite [16]. However, within the region C–D the hardness was increased from 360 Hv at point C to 500–525 Hv at point D, bottom of the substrate material, slightly lower than the substrate hardness of 580 Hv. The result suggests that the substrate material was overheated, causing an overextended tempering zone. The overheating can be avoided by increasing the substrate thickness and further reducing the heat input. Fig. 6a shows the clad hardness profile of the AerMet 100 asclad was approximately 560–580 Hv, which was 20% softer than the AISI 4340 as-clad. The difference in hardness of the two clads is attributed to the variation of composition, especially the reduced carbon content. Fig. 6b also shows the hardness of AISI 4340 as-clad þPCHT, where the hardness was consistent across the clad, HAZ, and overtempered zone, at approximately 630 Hv, except for the clad layer, where the average hardness is about 580 Hv. The reason is mainly due to a controlled heat treatment applied across the material. The average hardness was also 10% higher than the original substrate. Fig. 6b also shows the hardness profile of AerMet 100 asclad þPCHT was consistently 630 Hv across the HAZ and overtempered zone. The average hardness of the AerMet 100 clad layer after PCHT was 550 Hv. Fig. 7 shows the EDAXs composition analysis for the two clads and substrate. Fig. 7a shows the alloying elements of chromium and nickel were 0.35% and 1.0% higher, respectively, in the AISI 4340 clad, compared to the measured composition of AISI 4340 substrate and the composition provided by the supplier (Table 1). A similar result was observed in Direct Metal Deposition application of AISI 4340 steel [13]. The composition of the clad is expected to be different since a percentage of iron, alloying elements, and possibly carbon will evaporate due to extremely high heating temperatures during the cladding process [13]. Fig. 7b shows the composition analysis of AerMet 100 clad with a slightly higher nickel and cobalt, as compared to the composition provided by the supplier (Table 1). 3.3. Residual stresses Residual stresses play a very important role in structural components because these stresses can affect the fatigue and fracture behaviour. High tensile residual stresses can promote premature failure (sometimes internally) and compressive residual stresses can delay failure, resisting crack opening [19]. The residual stress as a function of depth was measured, as shown in Fig. 8. Since the substrate is rectangular and bi-axial in the prepared samples, the stresses in the two in-plane directions are similar and therefore only one principal stress result is shown. Fig. 8a shows, for both the AISI 4340 and AerMet 100 as-clads, the residual stress variations were similar. A compressive residual stress between 200 and 400 MPa is produced in the clad region and a sharp translation into tension residual stress occurs at the interface between the clad and the HAZ. The compressive residual stress produced in the clad layer is caused by thermal shrinkage,

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Fig. 7. EDAXs composition analysis: (a) AISI 4340 clad and substrate and (b) AerMet 100 clad.

thermal expansion–contraction, and displacive phase transformation of austenite to martensite [20,21]. A maximum tensile residual stress of 630 MPa (40% of the yield strength) is observed in the HAZ at approximately 0.5–1.0 mm below the clad interface. The residual stress decreases linearly with increasing distance for the remainder of the substrate, as would be expected as the material is further from the expanded clad material causing the stress. The tensile residual stress is a result of the thermal gradient created by the laser during laser cladding where during cooling the colder material constrains the hotter material, thus generating a tensile residual stress [21]. It should be acknowledged that a compressive stress due to the austenite martensite transformation would have reduced the magnitude of the tensile residual stress formed due to the previous mechanism [22]. It is a general observation that the maximum tensile residual stress, caused by laser cladding in steel, is in the HAZ adjacent to the clad interface [21,23]. Fig. 8b shows the importance and effectiveness of PCHT to minimise the residual stresses. A number of other studies have shown post-heat treatment to be an efficient method to reduce residual stresses [19,24]. The minimised and remaining residual stresses after PCHT are due to re-austenitising of the material and subsequent slow cooling. 3.4. Tensile properties Fig. 9 shows the comparison of the tensile behaviour with the tensile properties summarised in Table 3. The tensile properties of the substrate material are in agreement with a typical quench and

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Fig. 8. Longitudinal residual stress profile, measured relative to the clad interface: (a) as-clad condition and (b) PCHT condition.

220 1C temper AISI 4340 steel [4]. The substrate with 0.7 mm grind-out represents a notched specimen where the cross-section area used was of the original substrate rather than the reduced area. Therefore the tensile properties were reduced by 10–30%. Fig. 9a shows both the yield and the ultimate tensile strength (UTS) of the AISI 4340 as-clad were degraded by 40% compared to the 0.7 mm grind-out specimen (refer to Table 3). The failure of the AISI 4340 as-clad was very brittle, resulting in very little plastic deformation and the ductility (elongation) was reduced significantly by a factor of 20 compared to the 0.7 mm grind-out specimen. The brittleness is due to the very high hardness in the clad and HAZ from the predominance of untempered martensite, which exhibits low ductility and toughness. The failure was also pre-mature – attributed to the inhomogeneity of the material creating many local stress concentrators. The AerMet 100 as-clad shows that changing the composition of the clad, with a lower carbon and a higher nickel-cobalt content, can improve the tensile properties relative to those of the AISI 4340 as-clad. The PCHT can further improve tensile properties relative to those of the as-clad due to control of the microstructure, as shown in Fig. 9b. 3.5. Fatigue life Fig. 10 shows a comparison of the fatigue life curves. Again, the substrate with 0.7 mm grind-out was treated as a notched specimen, and as expected, failed below the substrate. Fig. 10a shows, for all stress levels tested, the fatigue life of the AISI 4340 as-clad was poor, especially at high cyclic loads (1200 MPa) where the fatigue life was reduced significantly to 221 cycles, compared to

Fig. 9. Stress–strain curve of the axial tensile test. The end points represent the point of fracture: (a) as-clad condition and (b) PCHT condition.

Table 3 Comparison of the tensile properties. 0.2% Yield (MPa) UTS (MPa) Substrate Substrate w/0.7 mm grind-out AISI 4340 as-clad AerMet 100 as-clad AISI 4340 as-clad þPCHT AerMet 100 as-clad þPCHT

1551.0 7 10 1442.5 7 60 1167.3 7 150 14727 20 1287.3 7 13 1298.8 7 8

Elongation (%)

1914.0 7 7 167 0.9 1769.8 7 25 11.17 1.9 1167.3 7 150 0.6 7 0.1 1616.3 7 80 3.0 7 0.6 2063.07 9 5.6 7 0.1 2110.5 7 21 7.2 7 0.5

the grind-out substrate of 8461 cycles. Similarly, at low cyclic loads (600 MPa), the fatigue life was also reduced significantly to 16,222 cycles, compared to the grind-out substrate of 87,386 cycles. Fig. 10a also shows, for all stress levels, the fatigue life of AerMet 100 as-clad improved by a factor of 10 compared to the AISI 4340 as-clad. However, at 1200 MPa cyclic load, the fatigue life of AerMet 100 as-clad (2605 cycles) was still reduced, compared to the grind-out substrate of 8461 cycles. However, at an 800 MPa cyclic load, the fatigue life of AerMet 100 as-clad was similar to the substrate grind-out. Fig. 10b shows, for all stress levels, AerMet 100 as-clad þPCHT (12,851 cycles at 1200 MPa cyclic loads) can improve the fatigue life relative to the substrate grind-out (8461 cycles). However, the AISI 4340 as-clad þPCHT fatigue life (4577 cycles at 1200 MPa cyclic loads) was still below the substrate grind-out (8461 cycles), primarily due to pre-mature crack development from porosity. The results indicated that the PHT could significantly increase fatigue lives of the as-clad materials.

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mode in the over-tempered zone consists of MVC with transgranular cracks. Fig. 12a shows the fatigue fracture surface appearance of AerMet 100 as-clad. The fatigue crack initiation occurred at the top surface of the specimen, as shown in Fig. 12b. The fracture mode near the crack initiation site is TTS (Fig. 12c). Other regions of the clad show MVC with local regions of quasi-cleavage (Fig. 12d). Very little IG fracture is observed in the clad, which is the main reason for an increased fatigue life compared to the AISI 4340 as-clad. Since the same substrate material is used, the fracture modes in the HAZ were similar to the AISI 4340 as-clad (refer to Fig. 11d–f), which shows IG and quasi-cleavage. Fig. 13a shows the fatigue fracture surface appearance of AISI 4340 as-clad þPCHT specimen. For all specimens, the crack initiated at the porosity, as shown in Fig. 13b, which was the main factor for the decreased fatigue life in these specimens. The fracture mode was quasi-cleavage (Fig. 13c), which is expected for an AISI 4340 steel that has been quenched and tempered [4]. Further away from the initiation site (Fig. 13d), tear ridges were more prominent in the fracture surface with local quasi-cleavage, in which crack propagation might have been more rapid in these regions. Fig. 14a shows the fatigue fracture surface appearance of the AerMet 100 as-clad þPCHT. The fracture started at the top surface of the clad region, as shown in Fig. 14b. The fracture mode in these specimens was mainly quasi-cleavage (Fig. 12c) accompanied by deep secondary cracking. Similar to the AISI 4340 as-clad þPCHT specimen, tear ridges were more prominent in the regions away from the initiation site (Fig. 14d).

4. Discussion Fig. 10. Constant amplitude axial S–N fatigue curve using R¼ 0.1 and frequency of 10 Hz: (a) as-clad condition and (b) PCHT condition.

3.6. Fracture modes Fig. 11a shows the fracture surface appearance of AISI 4340 asclad. The crack propagation arrested at the clad and HAZ interfaces due to grain discontinuity and possibly different fracture toughness of each region. The fracture occurred at the top surface of the specimen (clad region), and propagated rapidly in the clad layer, then HAZ, as shown in Fig. 11b. Large porosity with a diameter of 100–200 mm was also observed, although the fracture still initiated from the surface of the clad. Fig. 11c shows the clad layer exhibited mixed modes of fracture, mainly quasi-cleavage and intergranular (IG) along the prior austenite columnar grain boundary with regions of tearing topology surface (TTS) and tear ridges (TR). The IG fracture along the prior austenite grain boundaries is caused by IG embrittlement, possibly through segregation of manganese sulphide (MnS) at the prior austenite grain boundary [16,25]. IG fracture is generally encountered in low alloy cast steels [16]. TTS is a result of micro-plastic tearing and often associated with overload and high cyclic fatigue stresses [26]. Quasi-cleavage is a result of tear ridges formed by micro-void coalescence (MVC) accompanied by cleavage (C) facets [25]. No fatigue striations were observed and any form of fatigue fracture might have occurred by TTS. Fig. 11d shows the fracture mode in the thumbnail crack region in the HAZ with a distinct transition from IG to quasicleavage (Fig. 11e). Chevron markings were clearly visible in the HAZ (left to right in Fig. 11b) and generally indicate brittle fracture zones and rapid crack propagation, which would be expected of a material with very high hardness (4700 Hv) and a peak tensile residual stress of 600 MPa in the HAZ. Fig. 11f shows the fracture

4.1. Fatigue behaviour The fatigue behaviour of laser clad material is complex as the formation of a clad specimen produces a ‘composite’ with four individual regions (clad, interface, HAZ, and over-tempered zone), compared to a heat-treated uniform microstructure of the substrate. From a repair certification perspective and homogeneity of the clad specimen, it is ideal to use the same clad material as the substrate. Although the composition of the AISI 4340 clad is relatively similar to the substrate, with a slight increase in chromium and nickel content (Fig. 7a), each region has distinct and unique microstructure and mechanical properties due to variations in heating and rapid cooling – resulting in complex stresses. The AISI 4340 as-clad condition, while the substrate successfully restored the cross-sectional area from the grind-out repair, displayed reduced tensile properties (Fig. 9a) and fatigue life (Fig. 10b), in comparison with the substrate grind-out specimens. The major reasons for this degradation are (1) high hardness and brittle properties in the clad and HAZ region, (2) high tensile residual stresses in the HAZ and (3) grain non-uniformity such as the columnar grains in the clad and coarsening of the prior austenite grain boundary in the HAZ (Figs. 4a and 5b) – promoting intergranular fracture. The combination of these three factors caused a significant reduction in fatigue life. The bond strength at the clad interface is also an issue and can degrade the tensile properties – which was the case for H13 tool steel clad on copper alloy substrate [27]. In their study, the fracture initiated from the clad interface. However in this study, no delamination was observed for all conditions and the failure initiated from the top surface of the clad – evident that the fusion bond strength was strong. A compressive residual stress is produced in the AISI 4340 clad layer (Fig. 8a), and

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Fig. 11. SEM images of AISI 4340 as-clad fatigue specimen failed at 221 cycles at 1200 MPa maximum cyclic load: (a) fracture surface appearance, (b) circled crack initiation site, (c) fracture mode in clad, (d) fracture modes in the HAZ, (e) fracture modes in the HAZ Chevron region and (f) fracture modes in the over-tempered zone. (IG ¼intergranular, TTS¼ tearing topology surface, TR¼ tear ridges, C ¼ cleavage facets, and MVC ¼ micro-void coalescence.)

this was expected to provide resistance to crack development. However, the fracture still initiated from the surface of the AISI 4340 clad (Fig. 11b); this is believed to be due primarily to the poor mechanical properties (high hardness, low toughness, and low capacity to work harden) of the clad layer. 4.2. Influence of microstructure The high hardness and brittle properties in the clad and HAZ are due to the formation of un-tempered martensite phase in these regions. Quench and temper AISI 4340 grade steel achieves strength through the phase transformation of austenite to martensite, and retains toughness at the expense of a slight loss in strength, through tempering of the martensite [4]. In laser cladding of AISI 4340 steel, during cooling of the deposited track,

austenite to martensite transformation occurs in the clad and HAZ due to the rapid cooling rates, which for laser cladding are usually between 103 and 105 1C/s [2,13]. At these cooling rates, the transformation of ferrite from austenite is suppressed and displasive process occurs, causing martensite to form from the austenite lattice, which is one of the hardest and most brittle phase of steel. Bhattacharya et al. [13] reported on the microstructure evolution in direct metal deposition application of high strength steel AISI 4340 steel powder. In their study, the AISI 4340 steel clad microstructure primarily consists of tempered lath martensite. Tempering occurred due to the re-heating caused by the subsequent overlapping tracks and hence significant tempering of the martensite occurs. In this study, it is also reasonable to assume that the re-heating, caused by the subsequent overlapping tracks, will lead to the

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Fig. 12. SEM images of AerMet 100 as-clad fatigue specimen failed at 2206 cycles at 1200 MPa maximum cyclic load: (a) fracture surface appearance, (b) circled crack initiation site, (c) fracture mode in the clad near the initiation site and (d) fracture mode in the crack propagation region of the clad. (IG ¼intergranular, TTS ¼tearing topology surface, TR¼ tear ridges, C ¼ cleavage facets, and MVC ¼ Micro-void Coalescence.)

tempering of the martensite, causing a reduced hardness. However, the hardness remained uniform in the AISI 4340 clad layer at approximately 650 Hv (Fig. 6a), which is similar to the hardness of untempered martensite of AISI 4340 after an oil quench heat treatment [4]. For tempering to occur, the previous track must cool down below the martensite start (Ms) temperature before the re-heating of the overlapping track. It seems likely that tempering on the overlapping track was suppressed as the austenite in the previous track was not cooled down below Ms temperature to form martensite – attributed to the progressive rise of the substrate temperature above the Ms temperature [28], which for AISI 4340 Ms ¼298 1C. The clad and HAZ still remained austenitic during the re-heating of the overlapping tracks, where austenite to martensite transformation occurred only after the completion of the cladding when the bulk material was cooled down. Thus, hard and brittle untempered martensite remained in the clad and HAZ, which consequently led to the decreased fatigue life and brittle failure modes of AISI 4340 as-clad. Increasing the substrate size would allow the substrate to retain its heat extraction capacity [28], causing the deposited track to cool down below the Ms temperature and promoting the tempering effect from the overlapping track – a preferred heat treatment for increasing ductility and toughness. 4.3. Influence of AerMet 100 powder AerMet 100 as-clad demonstrated a great increase of fatigue life, as compared to AISI 4340 as-clad. AerMet 100, known as secondary hardening steels, achieves high strength and toughness through the precipitation of M2C carbides in the lath martensite during age

tempering at 482 1C [29]. As discussed previously, the tempering effect of the overlapping clad track was suppressed. Hence, little or no secondary hardening takes place in the AerMet 100 clad. The microstructure formation of AerMet 100 clad was similar to AISI 4340 (Fig. 4a and b). The high nickel content in AerMet 100 would produce Fe–Ni lath martensite and promote ductility. The reduced clad hardness (Fig. 6a) was due to the lower carbon content of AerMet 100 (0.25 wt%), compared to AISI 4340 (0.4 wt %). It seems reasonable that reduced carbon content and increased nickel content in the AerMet 100 clad would improve the fatigue life by reducing the brittleness and increased toughness. This is evident in the fracture surface of AerMet 100 clad where the primary mode of fracture is TTS and quasi-cleavage (Fig. 12c and d), compared to the dominant IG fracture in the AISI 4340 clad (Fig. 11c). However, similar to the AISI 4340 clad, even with a compressive residual stress, the fracture still initiated from the surface of the AerMet 100 clad, indicating the hardness of 550 Hv is still too high and brittle. In future, using softer grade steels for the clad material, such as HY 180 (carbon content of 0.10 wt%) and AF 1410 (carbon content of 0.15 wt%), may provide further benefit on fatigue properties. Since the as-clad is a composite-like structure, as a consequence of the reduced strength in the softer grade steels, the clad will yield first. However, softer grade steels have very high elongation and a good capacity for work hardening; and with the assistance of the compressive stress generated in the clad (Fig. 8a), it would lead to a more uniform stress re-distribution into the HAZ and substrate. Using softer grade high strength steels as the clad layer will be investigated as a future work to further improve the fatigue properties of ultra-high strength steels.

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Fig. 13. SEM images of AISI 4340 as-clad þ PCHT fatigue specimen failed at 4316 cycles at 1200 MPa maximum cyclic load: (a) fracture surface appearance, (b) circled crack initiation site, (c) fracture modes near the initiation site and (d) fracture mode in the crack propagation region. (IG ¼intergranular, TTS¼ tearing topology surface, TR ¼tear ridges, C ¼ cleavage facets, aMVC ¼micro-void coalescence.)

4.4. Influence of PCHT The experimental results clearly demonstrate that PCHT significantly increases the fatigue resistance to crack development and propagation, showed by an increase in tensile properties and fatigue life relative to those without post-heat treatment (Figs. 9 and 10). It should be noted that even though the compressive residual stress in the clad layer was removed from PCHT (Fig. 6b), the fatigue life was still improved – this is believed to be due primarily to the recovered mechanical properties (especially ductility and toughness) from PCHT in the clad and HAZ region. For future work, it is important that the material properties in the substrate are not altered during PCHT. It seems likely that use of a post-repair heat treatment process – preferably localised rather than bulk, to avoid dimensional distortion and further damage to the substrate – would temper the clad and HAZ only, while also reducing the tensile residual stresses in the HAZ. The benefits of localised post-heat treatment – the substrate will be unaffected and it is unlikely an aeronautical component will be post-heat treated as a bulk material due to complications with certification and dimensional distortion. No information on localised post-cladding heat treatment is available in the open literature. Localised post-repair heat treatment will involve local heating from the laser beam onto the surface of the clad, and ideally, to achieve a desired tempering temperature through the clad and HAZ region, preferably 482 1C for AerMet 100. Controlled tempering at the appropriate temperature will benefit the AerMet 100 clad – allowing secondary

hardening effects of precipitation of M2C carbides and improving toughness while retaining the high strength.

5. Conclusion The following conclusions are drawn:

 The tensile properties and fatigue life of the AISI 4340 as-clad





are degraded compared to the substrate material. The degraded mechanical properties are associated with the very hard and brittle properties in the clad and HAZ region. Changing the clad layer to secondary hardening grade steel (AerMet 100) can improve both the tensile properties and fatigue life, as compared to those produced by the quench and temper grade steel powder (AISI 4340). Post-heat treatment demonstrated an increase in tensile properties and fatigue life relative to that without post-heat treatment.

The repair of ultra-high strength steels requires further development to restore the fatigue properties. Since the control of thermal cycling during the cladding process is limited, manipulating the powder composition offers opportunities for improving the tensile and fatigue properties, as demonstrated with AerMet 100. Also, it seems likely that use of a post-repair heat treatment process – preferably localised to avoid dimensional distortion – which would temper the clad and HAZ, while also reducing the residual stresses in the HAZ, provides the most promising path to

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Fig. 14. SEM images of AerMet 100 as-clad þ PCHT fatigue specimen failed at 12,851 cycles at 1200 MPa maximum cyclic load: (a) fracture surface appearance, (b) circled crack initiation site, (c) fracture mode near the initiation site and (d) fracture modes in the crack propagation region. (IG¼ intergranular, TTS¼ tearing topology surface, TR ¼tear ridges, C¼ cleavage facets, and MVC¼ micro-void coalescence.)

restore and possibly increase the fatigue life of damaged ultra-high strength steel components. Acknowledgements The authors wish to acknowledge the support of the Defence Materials Technology Centre (DMTC), which was established and is supported by the Australian Government’s Defence Future Capability Technology Centre (DFCTC) initiative. The authors would also wish to acknowledge the financial support of Directorate General Technical Airworthiness (DGTA). In addition, the authors are indebted to the Industirla Research Institute Swinburne (IRIS), especially Mr. Girish Thipperudrappa for his assistance during the laser cladding experiments. References [1] L. Sexton, S. Lavin, G. Byrne, A. Kennedy, J. Mater. Process. Technol. 122 (2002) 63–68. [2] R. Vilar, J. Laser Appl. 11 (1999) 64–79. [3] R.G. Bonora, H.J.C. Voorwald, M.O.H. Cioffi, G.S. Junior, L.F.V. Santos, Proc. Eng 2 (2010) 1617–1623. [4] W.S. Lee, T.T. Su, J. Mater. Process. Technol. 87 (1997) 198–206. [5] J.R. Kattus, Aerospace Structural Metals Handbook, first ed., Belfour Stulen Inc., Michigan, 1977. [6] B. Banerjee, Int. J. Solid Struct. 44 (2007) 834–859. [7] M. Fastow, M. Bamberger, N. Nir, M. Landkof, Mater. Sci. Technol. 6 (1990) 900–904. [8] R.L. McDaniels, S.A. White, K. Liaw, L. Chen, M.H. McCay, P.K. Liaw, Mater. Sci. Eng. A 485 (2008) 500–507.

[9] M. Brandt, S. Sun, N. Alam, P. Bendeich, A. Bishop, Int. Heat. Treat. Surf. Eng. 222 (2009) 105–114. [10] W. Wang, A.J. Pinkerton, L.M. Wee, L. Li, in: S. Hinduja, K.-C. Fan (Eds.), Proceedings of the 35th International MATADOR Conference, Springer, London, 2007, pp. 345–350. [11] Q. Liu, M. Janardhana, B. Hinton, M. Brandt, K. Sharp, Int. J. Struct. Integr. 2 (2011) 314–331. [12] S. Niederhauser, B. Karlsson, Mater. Sci. Technol. 19 (2003) 1611–1616. [13] S. Bhattacharya, G.P. Dinda, A.K. Dasgupta, J. Mazumder, Mater. Sci. Eng. A 528 (2011) 2309–2318. [14] S.D. Sun, Q. Liu, M. Brandt, M. Janardhana, G. Clark, 28th International Congress of the Aeronautical Sciences, Brisbane, 2012. [15] Aerospace Material Specification (AMS) 2759/2F, SAE International, 2010. [16] H.K.D.H. Bhadeshia, R.W.K. Honeycombe, Steels Microstructure and Properties, third ed., Butterworth-Heinemann, London, 2005. [17] D.S. Martin, F.G. Caballero, C. Capdevila, C.G. de-Andres, Mater. Trans. 45 (2004) 2797–2804. [18] R.K. Shiue, C. Chen, Metall. Trans. A 23 (1992) 163–170. [19] P. Bendeich, N. Alam, M. Brandt, D. Carr, K. Short, R. Blevins, C. Curfs, O. Kirstein, G. Atkinson, T. Holden, R. Rogge, Mater. Sci. Eng. A 437 (2006) 70–74. [20] M. Hidekazu, B. Miloslav, V. Adan, R. Sherif, D. Catrin, D. David, N. Kamran, Trans. Join. Weld. Res. Inst 37 (2008) 75–80. [21] J.Y. Chen, K. Conlon, L. Xue, R. Rogge, Mater. Sci. Eng. A 527 (2010) 7265–7273. [22] P.J. Withers, H.K.D.H. Bhadeshia, Mater. Sci. Technol. 17 (2001) 366–375. [23] A. Suárez, J.M. Amado, M.J. Tobar, A. Yáñez, E. Fraga, M.J. Peel, Surf. Coat. Technol. 204 (2010) 1983–1988. [24] D. Thibault, P. Bocher, M. Thomas, M. Gharghouri, M. Côté, Mater. Sci. Eng. A 527 (2010) 6205–6210. [25] A.F. Liu, Mechanics and Mechanisms of Fracture: An Introduction, first ed., ASM International, Ohio, 2005. [26] A.W. Thompson, J.C. Chesnutt, Metall. Trans. A 10 (1979) 1193–1196. [27] M.K. Imran, S.H. Masood, M. Brandt, S. Bhattacharya, J. Mazumder, Mater. Sci. Eng. A 528 (2011) 3342–3349. [28] L. Costa, R. Vilar, T. Reti, A.M. Deus, Acta Metall. 53 (2005) 3987–3999. [29] J.H. Graves, Master's thesis, Worcester Polytechnic Institute, Worcester, Massachusetts, United States, 1994.