Effects of microstructures on the fatigue crack growth behavior of laser additive manufactured ultrahigh-strength AerMet100 steel

Materials Science & Engineering A 721 (2018) 251–262

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Effects of microstructures on the fatigue crack growth behavior of laser additive manufactured ultrahigh-strength AerMet100 steel

T

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Xian-zhe Rana,b,c, Dong Liua,b,c, Jia Lia,b,c, , Hua-ming Wanga,b,c, Xu Chenga,b,c, Ji-kui Zhanga,b,d, Hai-bo Tanga,b,c, Xiao Liuc a

National Engineering Laboratory of Additive Manufacturing for Large Metallic Components, 37 Xueyuan Road, Beijing 100191, China Engineering Research Center of Ministry of Education on Laser Direct Manufacturing for Large Metallic Component, 37 Xueyuan Road, Beijing 100191, China c School of Materials Science and Engineering, Beihang University, 37 Xueyuan Road, Beijing 100191, China d Department of Aircraft Design, Beihang University, 37 Xueyuan Road, Beijing 100191, China b

A R T I C L E I N F O

A B S T R A C T

Keywords: Laser additive manufacturing Ultrahigh strength steel Microstructure texture Fatigue crack growth Fracture resistance

In order to evaluate the effects of microstructure characteristics on fatigue crack growth (FCG) resistance of laser additive manufactured (LAM) AerMet100 steel, microstructures and FCG behaviors (in Paris region) of as-deposited specimen and three types of tempered martensite specimens were examined. Results indicate as-deposited specimens of LAM AerMet100 steel have apparent texture characteristics of epitaxy unidirectional growth prior-austenite columnar grains and grain-interior inter-dendritic blocky retained austenite with [001] crystallographic orientation. And poor boundary cracking resistance of these texture characteristics along deposition direction mainly contributes to the FCG rate anisotropy of as-deposited specimens. After post-LAM heat treatments, the FCG resistance of all heat-treated specimens apparently improves with the fracture mode of transgranular cracking. With the increase of yield strength, the value of Paris coefficient C of the steel increases, but the value of Paris exponent m decreases. Compared to the poor dislocation slip resistance of bainite plates in as-deposited specimens, the improved dislocation slip resistance of martensite plates is mainly related to the strong dislocation pinning effect of fine dispersive rod-like coherent M2C carbides, resulting in the stronger FCG resistance of the heat-treated specimens. In the Paris region of low ΔK (< ~ 20 MPa m1/2), fatigue cracks mainly propagate along the bainite (or martensite) plate interfaces, and the FCG rate of the steel can be effectively decreased by containing higher contents of thick film-like retained/reverted austenite; with the increase of ΔK, besides propagating along the soft inter-plate film-like austenite, fatigue cracks can also directly pass through the harder bainite (or martensite) plates with the striations and secondary cracks observed on fracture surfaces; in the Paris region of high ΔK (> ~ 70 MPa m1/2), higher contents of retained/reverted austenite inversely accelerate the FCG rate of heat-treated LAM AerMet100 steel. In contrast, grain refinement has the little influence on the FCG rate (in most of Paris region) of the heat-treated specimens.

1. Introduction Good fatigue resistance is always desired to high trustworthiness and durability of fracture critical components of advanced aircrafts with minimum life-cycle cost [1,2]. AerMet100 steel, 23Co14Ni11Cr3Mo, is an important high Co-Ni secondary hardening ultrahighstrength steel (UHSS). With a slightly overaged treatment (at 482 °C for 5 h), the microstructures of AerMet100 steel mainly compose of unrecovered highly dislocated Fe-Ni martensitic packets/plates, plate-interior dispersive distribution of fine coherent M2C (M=Cr, Mo, Fe) carbides and inter-plate thin film-like retained/reverted austenite, which result in high yield strength, good fracture toughness and

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extremely excellent fatigue resistance and hydrogen assisted cracking resistance [3–6]. Thus, it has been used in high stressed components (such as landing gear sleeve and arresting hook) of current and next generation of naval aircrafts [5,7]. According to two-stage model [8], fatigue fracture of alloys includes crack initiation stage and subsequent propagation stage. In the conventionally forged metallic materials, fatigue micro-cracks are prone to competitively initiate in regions containing of non-metallic inclusions, surface persistent slip bands (PSB) and microstructure heterogeneities (e.g., grain boundaries or phase interface) [9–12]. When stress strength factor range (ΔK) in these crack initiation regions exceeds fatigue threshold (ΔKth), fatigue cracks propagate, and which finally shorten

Corresponding author at: National Engineering Laboratory of Additive Manufacturing for Large Metallic Components, 37 Xueyuan Road, Beijing 100191, China. E-mail address: [email protected] (J. Li).

https://doi.org/10.1016/j.msea.2018.02.088 Received 7 November 2017; Received in revised form 23 February 2018; Accepted 25 February 2018 Available online 27 February 2018 0921-5093/ © 2018 Elsevier B.V. All rights reserved.

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the service life of components. As a safe-life structure, an extremely low level of fatigue risk is strictly required for landing gear. Therefore, the stress-controlled fatigue, strain-controlled fatigue and fatigue crack growth (FCG) behaviors of AerMet100 steel were evaluated by many studies [5,13–17]. The FCG behavior and microstructure-related cracking resistance are important fatigue performance indexes for evaluating remnant life of metallic component with micro-cracks [10,18,19]. Lee [5] investigated the FCG behaviors of AerMet100 steel under constant amplitude loading in both inert and corrosive environment. With an increase of stress ratio, the fatigue threshold of the steel decreased, but near-threshold FCG rate increased. It was noted that the corrosive environment could also slow down the near-threshold FCG rate of the steel due to corrosive-product-induced crack closure. However, when FCG rate of the steel was above 10−6 in./cycle, the corrosive environment and stress ratio had little effect on FCG rate of the steel. Furthermore, the study of Newman et al. [13] indicated cryogenic environment had ignored influence on the fatigue threshold and FCG rate of AerMet100 steel. In comparison with the effects of mechanical and environment factors, the influence of microstructural characteristics on the FCG behavior of AerMet100 steel is rarely reported in above literatures [5,13]. Laser additive manufacturing (LAM) is one of the additive manufacturing technologies which can be used for producing near-shape, dense and large alloy parts with complicated shapes by melting and solidification inert-gas-carried pre-alloyed powders in a layer-by-layer deposition mode [20]. In comparison with the conventional forging process, LAM AerMet100 steel component has the advantages of no mold, reduced material wastes and post-deposition machining, shorter production cycle, lower production cost and unprecedented design flexibility, etc. [20,21]. Generally, as-deposited specimens of LAM AerMet100 steel have large-size prior-austenite columnar grains with interior unidirectional growth of core-dendrite solidification structure and complex room-temperature microstructures, resulting in low strength and some ductility anisotropy [22]. After proper post heat treatments, the LAM AetMet100 steel can achieve superior tensile mechanical properties and fracture toughness due to microstructure optimization (including grain refinement, decrease of retained austenite, dissolution of M3C carbides and dispersive precipitation of coherent M2C carbides) [3,22–25]. Considering the good static mechanical properties, a major concern for the large LAM AerMet100 steel component is the evaluation of microstructure-related cyclic fatigue performance. In our previous study [23], fatigue strength and crack initiation sites of heat-treated LAM AerMet100 steel were reported. However, the influences of both inhomogeneous columnar-grain asdeposited microstructures and optimally equiaxed-grain tempered microstructures on FCG behaviors of the LAM AerMet100 steel have not been investigated. As it was known, the change of microstructure constituents (such as grain size, secondary phase precipitates, retained austenite content, etc.) could contribute to different FCG resistance and cracking behavior of alloys [26–32]. Higher fatigue threshold and FCG resistance of alloys were reported to be associated with larger grain size and higher contents of stable retained austenite [29,33]. Additionally, secondary hardening have large influence on the FCG resistance and cracking behavior (in Paris region) of precipitation hardening steel [30]. In this paper, microstructure characteristics and FCG behaviors of as-deposited specimens and three types of heat-treated specimens were firstly examined; and FCG rate anisotropy of as-deposited specimens and microstructure-related FCG resistance (in Paris region) of heattreated specimens were further discussed.

Fig. 1. As-deposited thick plate of LAM AerMet100 steel fabricated by pre-alloyed powders.

rotation electrode atomization process with diameters from 75 to 250 µm (in Fig. 1). Chemical composition of the powders (wt%) was: 13.50Co, 11.26Ni, 3.00Cr, 1.25Mo, 0.22C, 0.022 Nb, 0.022 Si, < 0.005Mn, 0.011 Al, 0.0007 S, < 0.005 P, < 0.005Ti, 0.0041O, 0.0009N, with the balance of Fe. Detailed information of LAM AerMet100 steel process was reported in our previous study [22]. A thick-wall plate with dimensions of 240 mm × 45 mm× 370 mm was finally fabricated (in Fig. 1). A large bulk with the dimensions of 170 mm × 45 mm× 280 mm was firstly cut from middle-upper to bottom of the as-deposited thickwall plate and followed by a preliminary heat treatment in hot isostatic pressing (HIP) furnace. After HIP, three bulks with the dimensions of 70 mm × 45 mm× 80 mm were further cut from the HIPed bulk for subsequent heat treatments. The heat treatment procedures for achieving different tempered martensitic microstructures are shown in Table 1. Finally, three kinds of heat-treated microstructures were designed as follows: (1) coarse grain tempered martensite microstructure (CG-TM); (2) fine grain tempered martensite microstructure (FG-TM); (3) fine grain tempered martensite microstructure with high contents of retained austenite (FG-TM-HRA). And tensile mechanical properties of the specimens with the above microstructures are listed in Table 2. 2.2. Fatigue crack growth test Two sizes of compact tensile specimens were used for FCG tests. Considering ductility anisotropy of as-deposited microstructures [22] and limited width (45 mm) of thick-wall plate, the smaller compact tensile specimens (Fig. 2b) with both L-T direction and ST-L direction (in Fig. 2a) were prepared from the middle-upper part of as-deposited thick plate. In contrast, the larger compact tensile specimens (Fig. 2c) with only L-T direction were prepared from three types of heat-treated bulks. It should be noted specimen size do not influence the comparison between the relationships (da/dN vs ΔK) of as-deposited specimens and those of heat-treated specimens. The FCG tests were applied under constant amplitude loading at room temperature in air using a MTS 880 mechanical testing machine with a 10 kN sensor according to ASTM E647-15 standard [35]. Prior to each test, pre-cracks with the diameter of 0.15 mm and length of 1 mm were firstly made at the notch tip of specimens by electrical discharge machining (EDM); in order to achieve sharp initial crack-tip shape, fresh fatigue pre-cracks with length of 1 mm for as-deposited specimens and 1.5 mm for heat-treated specimens were further made at a frequency of 20 Hz and stress ratio of 0.1 with the maximum force of

2. Experimental 2.1. Microstructure preparation AerMet100 steel pre-alloyed powders were fabricated via plasma 252

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Table 1 Designed post-LAM heat treatment processes for achieving different tempered martensitic microstructures. No.

Alloy designation

Preliminary heat treatment

Subsequent heat treatments

1

CG-TM

1300 °C, 135 MPa, 5 h,furnace cooling (FC)

2

FG-TM

3

FG-TM-HRA

Special heat treatment−1 + 885 °C, 1 h, oil cooling (OC) + cryogenic treatment in liquid nitrogen, 2 h, air warming + 482 °C, 5 h, air cooling (AC) Special heat treatment-2 + 885 °C, 1 h, OC + cryogenic treatment in liquid nitrogen, 2 h, air warming + 482 °C, 5 h, AC special heat treatment-2 + 885 °C, 1 h, OC + 482 °C, 5 h, AC

surface were calculated with an assumed critical resolved shear stress (CRSS) of 0.2 for all matrix microstructures. In addition, alloy carbides and dislocation sub-structure were examined using a Tecnai G2 F30 STWIN field-emission high resolution transmission electrode microscope (TEM) operated at 300 kV and equipped with an EDAX energy-dispersive X-ray spectroscopy (EDS) system. And TEM thin slice specimens with the diameter of 3 mm taken from the tested FCG specimens were twin-jet electro-polishing using a mixed solution of 10% perchloric acid and 90% ethanol at − 15 °C and 200 mA.

1.5 kN. After that, FCG tests were carried out at a frequency of 20 Hz and stress ratio of 0.06 with the maximum force of 0.6 kN for as-deposited specimens and 1.2 kN for heat-treated specimens. The crack mouth open displacements for computing the compliance were measured using a 5 mm length crack opening displacement (COD) gauge fixed on integral knife edges machined across the crack mouth. The FCG rates (da/dN) were expressed as a function of apparent stress-intensity factor range (ΔK), which was calculated from the expressions based on linear elastic stress analysis [35]. And the FCG rates in the steady-state (Paris) region of the tested specimens can be characterized by Paris law.

da/ dN = C (∆K )m

(1)

3. Results

where C and m are the materials constants. When axial displacements were over the values of displacement detector limits, FCG tests of some specimens were automatic stopped due to fluctuation sensitivity limit of COD gauge. After tests, these unbroken specimens were completely separated by moving cross-heads of mechanical testing machine.

3.1. Microstructure characteristics In as-deposited LAM AerMet100 steel, large-size prior-austenite columnar grains (approximate 150–600 µm) with grain-interior epitaxy growth (about 18–38 µm) cellular solidification structure are distribution along the deposition direction (Fig. 3). And the microstructures of the as-deposited steel mainly consist of grain boundary allotriomorphic ferrite (GBA), grain interior irregular proeutectoid ferrite, lots of platelike upper bainite, needle-like lower bainite and textured retained austenite [22]. And upper bainite packets are always observed as the feather-like morphologies, which have the average size of about 60 µm (in Table 3). EBSD results (in Fig. 4b) show that blocky retained austenite phases are distribution in the inter-dendritic regions along [001] crystallographic direction (in Fig. 4). And misorientation angle of the blocky retained austenite interfaces is about 45° (Fig. 4c and d). Furthermore, color-coded orientation map of [001] inverse pole figure (IPF) (in Fig. 4a) shows the blocky retained austenite in neighbored grains have different crystallographic planes, suggesting a tilt and twist boundary between two grains. After post-LAM heat treatments, prior-austenite grains of LAM AerMet100 steel are significantly equiaxed and refined with decreased packet size (in Fig. 5 and Table 3). Although the average prior-austenite grain sizes of FG-TM and FG-TM-HRA specimens are much smaller than those of CG-TM specimen, the average martensitic packet sizes of the former specimens are quite close to that of the latter one (in Fig. 5 and Table 3). The results of misorientation angle distribution (in Table 3) suggest that high-angle boundaries (> 15°) take the predominant fraction (> 85%) for both as-deposited and all heat-treated specimens. In general, with a decrease of retained/reverted austenite, the number of high-angle austenite boundaries of LAM AerMet100 steel decrease.

2.3. Microstructure characterization and fractography Specimens with the sizes of 15 mm × 15 mm× 15 mm were sectioned from the middle parts of as-deposited plate and different heattreated bulks separately. Then, the XOY and XOZ plane of the as-deposited specimens and XOY plane of various heat-treated specimens were ground, mechanically polished, etched using a mixture of 4% Nital and saturated picric acid water solution and investigated using Leika-DM 4000 optical microscopy (OM) and JEOL JSM-6010LA scanning electrode microscopy (SEM). The average sizes of prior-austenite grains were measured by a linear intercept method [36]. For analyzing crystallographic orientations and boundary characteristics of microstructures, the ground specimens were electro-polished in a mixed solution of 4% perchloric acid and 96% ethanol, and subsequently examined by JSM-7001F SEM with Pegasus XM2 electron backscatter diffraction (EBSD) with a resolution of 0.15 µm. In addition, matrix phase compositions of different microstructures were identified using X-ray diffraction (XRD) with Cu Kα, scanning rates 0.5°/min, voltage 40 kV and current 200 mA. And retained/reverted austenite contents of matrix microstructures were quantitatively analyzed with the structure refinement by Rietveld method with a software of GSAS-II [37]. After FCG tests, fracture surfaces of specimens were directly examined using JEOL JSM-6010LA SEM. And fracture sub-surface of specimens was further cut, vibratory polishing and characterization by the EBSD with a resolution of 15 nm. All the EBSD results were analyzed by the software of OIM Analysis 6.2. Taylor factors of fracture subTable 2 Tensile mechanical properties of LAM AerMet100 steel. Alloy state designation

Elastic modulus (E),GPa

UTS(σ b ),MPa

YS(Rp0 . 2 ),MPa

EL(δ) , %

RA(ψ ),%

FG-TM FG-TM-HRA CG-TM As-deposited [22] Forged [4,6,15,25,34]

188 ± 1 173 ± 3 185 ± 1 – 190

2003 ± 11 2027 ± 5 1977 ± 13 1583 ± 38 1965–2020

1827 ± 13 1610 ± 20 1767 ± 4 1062 ± 10 1724–1830

12.7 ± 0.3 13.3 ± 0.2 12.5 ± 0.3 12.3 ± 1.5 14

67.2 ± 3.2% 62.4 ± 1.3% 56.3 ± 1.6% 36.4 ± 8.3% 61–65%

Note: The principle axis of tensile specimens was parallel to the deposition direction. And the detail information of tensile tests was described in literature [22].

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Fig. 2. (a) Schematic diagram showing fatigue crack growth specimens with different notch directions; geometries and dimension (mm) of as-deposited specimens (b) and heat-treated specimens (c).

of heat-treated LAM AerMet100 steel is sensitive to the change of retained/reverted austenite contents but rarely influenced by grain refinement. Moreover, the effects of retained/reverted austenite contents on FCG rate of LAM AerMet100 steel are closely related to the value of ΔK. At a low ΔK, higher retained/reverted austenite contents suggest a beneficial effect on the decreased FCG rate of LAM AerMet100 steel. For example, at a ΔK of 14 MPa m1/2, the FCG rate of FG-TM-HRA specimen is only about 63% of that of FG-TM specimen; in contrast, at a ΔK of 30 MPa m1/2, the former one increases into about 82% of the latter one. However, at a high ΔK (＞~ 70 MPa m1/2), higher retained/reverted austenite content inversely accelerate the FCG rate of LAM AerMet100 steel. Similarly, at the high ΔK, the FCG rate of CG-TM specimen becomes higher than that of FG-TM specimen. Both of these results further suggest the higher retained/reverted austenite content and the coarser prior-austenite grain are not beneficial to the improvement of FCG resistance of LAM AerMet100 steel at the high ΔK. In addition, it is noted that the values of Paris coefficient C and exponent m are strongly influenced by yield strength of LAM AerMet100 steel. And the increase of yield strength could increase the value of Paris coefficient C of the steel but decreases the value of Paris exponent m (in Tables 2, 4 and 5). In the Paris region of high ΔK (＞~ 70 MPa m1/2), the specimens of LAM AerMet100 steel with higher yield

3.2. Fatigue crack growth rate Fig. 6 and Table 4 is the FCG relationship (da/dN vs ΔK) and corresponding results of as-deposited specimens. In the Paris region, the FCG rate of AD-ST-L specimen is higher than that of AD-L-T specimen. Compared to AD-L-T specimen, with the increase of ΔK, the AD-ST-L specimen has a higher increment of FCG rate, resulting in a higher value of Paris exponent m. Furthermore, at a ΔK of about 45 MPa m1/2, fatigue crack in AD-ST-L specimen is prone to become fast propagation (stageⅢ). Above results suggest FCG rate of as-deposited LAM AerMet100 steel is anisotropy with poor cracking resistance along deposition direction. Fig. 7 and Table 5 is FCG relationship (da/dN vs ΔK) and corresponding results of heat-treated specimens separately. After post-LAM heat treatments, the heat-treated specimens of LAM AerMet100 steel have much larger Paris region (stage-II), lower values of Paris exponent m and higher values of Paris coefficient C. These results indicate postLAM heat treatments can apparently improve FCG resistance of LAM AerMet100 steel at a higher ΔK. In Paris regions of less than ~ 70 MPa m1/2, FCG rate of FG-TM specimen is similar to that of CG-TM specimen, and which is apparently higher than that of FG-TM-HRA specimen. These results indicate the FCG rate (in most of Paris region) 254

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Fig. 3. As-deposited microstructures of LAM AerMet100 steel: low magnification optical microscope images of the XOZ plane (a) and the XOY plane (b) showing large-size prior-austenite columnar grains and grain-interior epitaxy growth cellular solidification structure; high magnification SEM images showing the feather-like upper bainite packets and plates in the XOZ plane (c) and the XOY plane (d).

20 MPa m1/2, the interface cracking of blocky retained austenite in inter-dendritic regions can be noticed in fracture surface of AD-ST-L specimen. In contrast, at the similar low ΔK, prior-austenite columnar grain boundary is not the preferential crack propagation path. And fatigue cracks always pass through prior-austenite columnar grain boundary by cross-slip (in Fig. 9b). Until the value of ΔK is excess ~ 30 MPa m1/2, the intergranular cracking of as-deposited specimen is appearance (in Fig. 9d). After post-LAM heat treatment, all of the heat-treated specimens are failure by transgranular cracking (in Fig. 10). FCG paths of these heattreated specimens are similar, which change with the increase of ΔK. In the Paris region of low ΔK, fatigue cracks propagate along martensitic packet/plate boundaries, resulting in appearance of granular morphologies (in Fig. 10b and d). At a ΔK of ~ 32 MPa m1/2, narrow fatigue striations are observed in fracture surfaces of heat-treated specimens. Simultaneously, besides boundary cracking of the martensitic packet/plate, the propagation of fatigue cracks also directly go through martensitic plates. With increase of ΔK, the width of fatigue striation and the number of secondary cracks gradually increase. In the region of high ΔK (＞~ 65 MPa m1/2), a few of dimples can be also observed (in Fig. 10e and f). At a ΔK of ~ 20 MPa m1/2, there are apparently less secondary cracks in fracture surfaces of heat-treated specimens than those in asdeposited one (in Fig. 8d and 10d). Fig. 11 further shows fracture sub-

strength are prone to have a lower FCG rate (in Table 2, Figs. 6 and 7). Therefore, it can be agreed that the strengthening effect can improve FCG resistance of LAM AerMet100 steel in the Paris region of high ΔK. 3.3. Fracture surface morphologies Fig. 8 shows typical transgranular fracture surface morphologies of FCG specimens of as-deposited LAM AerMet100 steel. In Paris region of low ΔK (Fig. 8a and b), fatigue cracks of as-deposited specimens are mainly propagation along bainite plate interfaces. With the increase of ΔK, a plenty of secondary cracks are observed not only at the bainitic plate interfaces but also in the interiors of bainite plates (Fig. 8c and d), indicating a poor transgranular cracking resistance of bainite plates. At a ΔK of ~ 32 MPa m1/2, narrow fatigue striations are observed in fracture surfaces of as-deposited specimens (Fig. 8f), suggesting a plastic blunting and re-sharpening process of fatigue crack propagation. Additionally, severe secondary cracks, paralleling to the narrow fatigue striations and perpendicular to main crack propagation direction, are also observed. Besides transgranular cracking of bainite plate and its interfaces, transgranular cracking of blocky retained austenite interfaces in interdendritic regions and intergranular cracking of prior-austenite columnar grain boundaries can be also observed in fracture surfaces of asdeposited specimens (in Fig. 9). Even though at a ΔK of less than

Table 3 Prior-austenite grain size, packet size, retained/reverted austenite contents and boundary characteristics of LAM AerMet100 steel obtained by OM, XRD and EBSD. Microstructure

As-deposited CG-TM FG-TM FG-TM-HRA a

Sectioned plane

XOY XOZ XOY XOY XOY

The average prioraustenite grain size, μm

The average packet width, μm

248 ± 102

57 ± 27

72 ± 42 18 ± 6 18 ± 6

16 ± 3 12 ± 3 12 ± 3

Retained/reverted austenite contenta, wt%

~ ~ ~ ~ ~

9.3 8.6 1.1 1.5 2.9

Low angle boundary fraction (> 5 ° and < 15 °), %

High angle boundary fraction (> 15°), %

γ

α

γ

α

0.37 0.24 0.29 0.25 0.22

1.41 2.95 5.17 4.25 4.33

31.60 20.94 12.76 15.89 18.06

59.76 65.14 70.23 69.84 66.95

Compared to the smaller tested area of EBSD, the larger tested area of XRD can better reflect the average values of retained/reverted austenite in specimens.

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Fig. 4. EBSD results of as-deposited LAM AerMet100 steel. (a) Color-coded [001] inverse pole figure (IPF) map and (b) the corresponding phase composition and retained austenite distributions of XOZ plane; (c)(d) the interfaces of blocky retained austenite showing misorientation angle of ~ 45°; (e) matrix phase composition and retained austenite distribution of XOY plane and (f) [001] IPF of blocky retained austenite. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

the martensite plate interfaces. This result may indicate critical resolved shear stress (CRSS) of tempered martensite plate is actually much higher than that of bainite plate.

surface morphologies color-coded by Taylor factor distribution maps of two representative FCG specimens of LAM AerMet100 steel at the ΔK of ~ 20 MPa m1/2. The major difference of these fracture sub-surface morphologies (in Fig. 11) is focused on the FCG path of the steel in the regions of high Taylor factor. For as-deposited specimen, in the region of high Taylor factor, cracks not only propagate along the bainite plate interfaces but also directly pass through the interior of bainite plates. However, for the heat-treated specimens, cracks only propagate along 256

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Fig. 5. Prior-austenite grain morphologies and EBSD color-coded [001] inverse pole figure maps of heat-treated LAM AerMet100 steel: (a)(c) CG-TM; (b)(d) both of FG-TM and FG-TMHRA. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

Fig. 6. Fatigue crack growth relationships (da/dN vs ΔK) of as-deposited LAM AerMet100 steel.

Fig. 7. Fatigue crack growth relationships (da/dN vs ΔK) of heat-treated LAM AerMet100 steel.

Table 4 Fatigue crack growth rate and Paris law parameters of as-deposited LAM AerMet100 steel. Specimen designation

AD-ST-L AD-L-T

Fatigue crack growth rate (mm/cycle)

Paris law parameters

ΔK = 14 MPa m1/2

ΔK = 20 MPa m1/2

ΔK = 30 MPa m1/2

ΔK = 50 MPa m1/2

Coefficient, C

Exponent, m

2.40E−5 1.35E−5

7.23E−5 4.06E−5

2.63E−4 1.53E−4

1.05E−3 3.75E−4

5.38E−9 2.21E−8

3.15 2.53

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Table 5 Fatigue crack growth rate and Paris law parameters of heat-treated LAM AerMet100 steel. Specimen designation

CG-TM FG-TM FG-TM-HRA

Fatigue crack growth rate (mm/cycle)

Paris law parameters

ΔK = 14 MPa m1/2

ΔK = 20 MPa m1/2

ΔK = 30 MPa m1/2

ΔK = 50 MPa m1/2

ΔK = 80 MPa m1/2

Coefficient, C

Exponent, m

2.84E−5 2.90E−5 1.85E−5

7.05E−5 9.31E−5 4.10E−5

1.67E−4 1.68E−4 1.38E−4

4.67E−4 4.93E−4 3.15E−4

1.41E−3 1.29E−3 1.42E−3

8.44E−8 1.16E−7 4.22E−8

2.22 2.14 2.33

4. Discussions

characteristics of the steel, and the propagation of fatigue cracks is mainly influenced by the other microstructure characteristics (e.g., matrix phase orientation, types and distribution of precipitates, etc.) [30,32,38,39]; on the other hand, inhomogeneous as-deposited microstructures may contribute to rough-induced crack closure in both of near-threshold region and initial Paris region, which can greatly slow down the initial FCG rate of the steel [40]. With the increase of ΔK, compared to prior-austenite columnar grain boundaries, the interfacial areas of blocky retained austenite are more sensitive to the increment of ΔK (in Fig. 9c and d). When some part of main crack starts to propagate along the phase interfaces of blocky retained austenite, FCG rate of the specimen is accelerated. As a result, the AD-ST-L specimen suggests the higher FCG rate and Paris exponent m than those of AD-L-T specimen

4.1. Effects of microstructure texture on the FCG rate anisotropy of asdeposited LAM AerMet100 steel Above results (Figs. 6, 9 and Table 4) show that FCG rate of asdeposited LAM AerMet100 steel is anisotropy due to microstructure texture characteristics of both prior-austenite columnar grains and grain-interior inter-dendritic blocky retained austenite. However, at a low value of ΔK, FCG rate anisotropy of as-deposited microstructure is normally not significance for two reasons. On the one hand, the maximum stress intensity factor (Kmax) at crack-tip plastic zone is lower than the boundary/interface cracking resistance of texture

Fig. 8. Typical transgranular fracture surface morphologies of FCG specimens of as-deposited LAM AerMet100 steel in Paris region of different ΔK. (a)(b) ~ 14 MPa m1/2; (c)(d) ~ 20 MPa m1/2; (e)(f) ~ 32 MPa m1/2.

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Fig. 9. Three types of typical cracking paths for FCG specimens of as-deposited LAM AerMet100 steel in the XOZ plane: (a)(b) transgranular cracking of bainite plate and its interface; (c) transgranular cracking of inter-dendritic blocky retained austenite interface; (d) intergranular cracking of prior-austenite columnar grain boundary.

(in Table 4). In Paris region of higher ΔK, when stress concentration at crack-tip plastic zone exceeds the fracture resistance of grain boundary allotriomorphic ferrite (GBA), the intergranular cracking of prior-austenite columnar grains can further accelerate the FCG rate of the specimen (in Fig. 9d and literature [22]). And if intergranular cracking dominates the FCG behavior, the FCG specimen of the steel will be rapidly failure within a short time. In contrast, when crack growth direction is perpendicular to the deposition direction, fatigue cracks cannot rapidly propagate along these sensitive boundaries, the growth rate of which slows down. In addition, it is possibility that blocky retained austenite has some positive effect on reducing the FCG rate of AD-L-T specimen because strain-induced martensite transformation can consume the driving force for fatigue crack propagation [39]. Finally, it should be pointed out the poor boundary cracking resistance of microstructure texture characteristics mainly contributed to the FCG rate anisotropy of as-deposited LAM AerMet100 steel.

bainite(or martensite) plate. During FCG, the plastic deformation and martensitic transformation of inter-plate film-like retained austenite can contribute to the increase of FCG resistance of the steel, which blunt the crack tip and absorb the energy that could be used for fatigue crack propagation [41]. Compared to film-like retained austenite with the thickness of ~ 13 nm in as-deposited specimen (in Fig. 12a and b), after post-LAM heat treatments, film-like retained/reverted austenite in heattreated specimens have a thinner thickness of ~ 3–8 nm [3,25]. After finite element method simulation of crack propagation interaction with film-like austenite, Wang et al. [25] proposed the presence of film-like austenite could largely reduce stress concentration of crack-tip and decrease the crack growth rate of the steel by changing the direction of crack propagation. Moreover, additional stress concentration beyond crack-tip could still form behind the austenite layer until the thickness of film-like austenite increased into about 10–15 nm. At a low value of ΔK (≤ 20 MPa m1/2), fatigue crack is prone to propagate along the bainite (or martensite) plate interfaces, and a high contents of interplate film-like austenite with a thickness of 10–15 nm is beneficial to slow down the initial FCG rate of the steel. Therefore, compared to the initial FCG rate of FG-TM specimen and CG-TM specimen, the initial FCG rate of FG-TM-HRA specimen and AD-L-T specimen is lower (in Tables 4 and 5). In the Paris region of higher ΔK, besides propagating along the softer inter-plate film-like austenite, fatigue cracks can also directly pass through the harder bainite (or martensite) plates with the fracture characteristic of striations. From the view of crack-tip dislocation emission, the plastic blunting and re-sharpening process of crack extension is strongly related to dislocation slip resistance of matrix [38]. And the interaction between alloy carbides and dislocations is key factor that influences the resistance of dislocation slip in matrix. Fig. 12 shows alloy carbides and dislocation distribution characteristics of LAM AerMet100 steel after FCG tests. In as-deposited specimen, substantial needle-like M3C carbides with length of 113–134 nm and diameter of 13–17 nm are distribution in bainite plates as the major precipitates, which contribute to the poor dislocation slip resistance for the bainite plate. Under cyclic loading, a plenty of slip bands are formation in

4.2. Effects of microstructure characteristics on the FCG resistance in Paris region The results of FCG rate, Paris exponent m and fracture surface morphologies indicate FCG resistance of LAM AerMet100 steel is influenced by the changes of microstructure characteristics and the corresponding yield strength. Compared to as-deposited specimens, all heat-treated specimens have the apparent improvement of FCG resistance with the lower Paris exponent m and the lower FCG rate in the Paris region of high ΔK due to microstructure optimization. First of all, with elimination of microstructure texture characteristics and formation of equiaxed grains, fast propagation paths of fatigue cracks are not presence in heat-treated specimens. And all of the heat-treated specimens become transgranular cracking. It has been mentioned that transgranular cracking paths of both asdeposited and heat-treated LAM AerMet100 steel change with the increase of ΔK (Figs. 8 and 10). At a low value of ΔK (≤ 20 MPa m1/2), fatigue cracks are prone to propagate along the bainite (or martensite) plate interfaces because inter-plate film-like austenite is softer than 259

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Fig. 10. Typical fracture surface morphologies of FCG specimens of heat-treated LAM AerMet100 steel in Paris region of different ΔK. (a)(b) ~ 11 MPa m1/2; (c)(d) ~ 20 MPa m1/2; (e)(f) ~ 65 MPa m1/2.

bainite plates, the motion of dislocations are prone to be impeded by the brittle needle-like M3C carbides. And severe dislocation pile-up in the interfaces of brittle M3C carbides can contribute to the formation of micro-cracks in bainite plates due to stress concentration (in Fig. 8d). This behavior greatly deteriorates the cracking resistance of bainite plate. In contrast, dislocations in tempered martensite plates of heattreated specimens are uniform distribution, and which are pinned by fine dispersive needle-like coherent M2C carbides. And these randomly dispersive M2C carbides have the lengths of 6–10 nm, the diameter of 1–2 nm and the spacing of 6–8 nm. The excellent yield strength and high work hardening capacity of heat-treated specimens [15,22] indicate that a strong strengthening effect can be achieved by the dislocation pinning of these fine dispersive M2C carbides. As a result, the movement of crack-tip emitted dislocations in the favorite slip plane becomes difficult due to this strengthening effect, and the dislocation slip resistance of matrix martensite plate increases. During propagating of fatigue cracks, low strength retained/reverted austenite can be regarded as one important microstructure factor which can slow down the FCG rate of the steel by the formation of secondary cracks. The secondary crack is essentially a result of the crack branching, which can release the triaxial stress conditions and blunt the crack tip, thus leading to higher resistance to the crack propagation [41]. However, at a high value of ΔK (＞~ 70 MPa m1/2), the

Fig. 11. Fracture sub-surface morphologies color-coded by Taylor factor distribution maps of two representative FCG specimens of LAM AerMet100 steel at a ΔK of ~ 20 MPa m1/2: (a) AD-L-T specimen; (b) FG-TM specimen. Note: The value of critical resolved shear stress (CRSS) used for Taylor factor calculation is assumed as default value of 0.2. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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Fig. 12. Alloy carbides and dislocation distribution characteristics of LAM AerMet100 steel after FCG tests obtained by TEM: (a)(b) needle-like M3C carbides interaction with slip bands in one bainite plate of as-deposited specimen; (c) high density of uniform distributed dislocations in one tempered martensite plate of heat-treated specimen; (d) dislocation pinned by one fine needle-like coherent M2C carbide.

as-deposited specimens. (2) After post-LAM heat treatments, microstructure texture characteristics are elimination, and prior-austenite grains of LAM AerMet100 steel become equiaxed with largely reduced number of high angle austenite boundaries. In comparison with fine grain size (18 ± 6 µm) and low contents of retained/reverted austenite (~ 1.6 wt%) for FG-TM specimen, the CG-TM specimen and the FGTM-HRA specimen has the larger average grain size (72 ± 42 µm) and higher austenite contents (~ 2.9 wt%) separately. Compared to as-deposited specimens, the FCG resistance of all heat-treated specimens apparently improves with the fracture mode of transgranular cracking. (3) In the Paris region of low ΔK (< ~ 20 MPa m1/2), fatigue cracks mainly propagate along the bainite (or martensite) plate interfaces, and the FCG rate of the steel can be effectively decreased by containing higher contents of thick film-like retained/reverted austenite; with the increase of ΔK, besides propagating along the soft inter-plate film-like austenite, fatigue cracks can also directly pass through the harder bainite(or martensite) plates with the striations and secondary cracks observed on fracture surface; in the Paris region of high ΔK (＞~70 MPa m1/2), higher contents of retained/ reverted austenite inversely accelerate the FCG rate of heat-treated LAM AerMet100 steel. (4) Compared to the poor dislocation slip resistance of bainite plates in as-deposited specimens, the improved dislocation slip resistance of martensite plates is mainly related to the strong dislocation pinning effect of fine dispersive rod-like coherent M2C carbides, resulting in

low strength retained/reverted austenite in large-size plastic zone beyond crack-tip is prone to be cracking, which can be regarded as the new-formed cracking paths. In the subsequent extension of crack-tip, these new-formed cracking paths can greatly decrease the resistance of crack propagation, and thus contribute to the apparent acceleration of FCG rate of the steel (in Fig. 7). Compared to the effect of austenite content, the influence of grain refinement on the FCG rate (in most of Paris region) of heat-treated LAM AerMet100 steel is not obvious. However, it is still possible that grain refinement can affect the nearthreshold FCG rate of the steel [42,43]. Anyway, these results can indicate both of alloy carbides and retained/reverted austenite are the major factors that influence the FCG resistance (in the Paris region) of LAM AerMet100 steel. 5. Conclusions The microstructure characteristics and fatigue crack growth behaviors (in Paris region) of both as-deposited specimen and three types of tempered martensite specimens of LAM AerMet100 steel at a stress ratio of 0.06 were examined. The main conclusions are as follows (1) As-deposited specimens have apparent microstructure texture characteristics of epitaxy unidirectional growth prior-austenite columnar grains and inter-dendritic blocky retained austenite with [001] crystallographic orientation. The poor boundary cracking resistance of the microstructure texture characteristics along deposition direction mainly contribute to the FCG rate anisotropy of 261

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the stronger FCG resistance of the heat-treated specimens. In addition, grain refinement has the little effect on the FCG resistance (in most of Paris region) of the heat-treated specimens.

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