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Correlation between the intergranular brittleness and precipitation reactions during isothermal aging of an Fe–Ni–Mn maraging steel
Materials Science and Engineering A 490 (2008) 105–112
Correlation between the intergranular brittleness and precipitation reactions during isothermal aging of an Fe–Ni–Mn maraging steel S. Hossein Nedjad a,∗ , M. Nili Ahmadabadi b , T. Furuhara c a b
Faculty of Materials Engineering, Sahand University of Technology, P.O. Box 51335-1996, Tabriz, Iran School of Metallurgy and Materials Engineering, University of Tehran, P.O. Box 14395-731, Tehran, Iran c Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan Received 27 August 2007; received in revised form 25 December 2007; accepted 24 January 2008
Abstract Evolution of the intergranular brittleness of an Fe–10Ni–7Mn (weight pct) maraging steel was correlated with precipitation reactions during isothermal aging at 753 K. Intergranular brittleness of the Fe–Ni–Mn steel raises after aging treatment which occurs catastrophically at zero tensile elongation in the underaged and peakaged steels. The intergranular failure is attributed to grain boundary weakening due to the formation of coarse grain boundary precipitates associated with solute-depleted precipitate-free zones during isothermal aging. Further, evidences of planar slip bands were found within the grains of a peakaged specimen loaded by tensile deformation. Those inhomogeneously deformed bands were identified to apply large strain localization in the soft precipitate-free zones at grain boundaries which is assumed to fascinate microcracks initiation at negligible macroscopic strains in the underaged and peakaged steels. During further aging, concurrent reactions including (i) overaging of matrix precipitates, (ii) spheroidization of grain boundary precipitates, (iii) growth of precipitate-free zone in width and (iv) diffusional transformation to austenite take place which increase tensile ductility after prolonged aging. © 2008 Elsevier B.V. All rights reserved. Keywords: Maraging steel; Intergranular brittleness; Grain boundary precipitation; Precipitate-free zone; Planar slip
1. Introduction Iron–nickel–manganese martensitic steels show substantial age hardening but suffer from poor ductility after aging [1,2]. The Fe–Ni–Mn maraging steels exhibit intragranular dimpled ductile fracture in the solution-annealed condition which turns, by short isothermal aging, to premature intergranular brittle fracture, passing along prior austenite grain boundaries (PAGBs) [3]. Hereafter, the PAGBs will be refereed to as grain boundaries. Analogous to the temper-embrittlement of low alloy steels, Squires and Wilson [3] first suggested that intergranular failure of an Fe–12Ni–6Mn (weight pct) maraging steel arises from segregation of manganese at grain boundaries during isothermal aging at 573–773 K. Hereafter, all chemical compositions will be given in weight pct. Then Feng et al. [4] showed by Auger electron spectroscopy (AES) that nickel and manganese segregate at grain boundaries of an
∗
Corresponding author. Tel.: +98 412 345 9449; fax: +98 412 344 4333. E-mail address: [email protected] (S. Hossein Nedjad).
0921-5093/$ – see front matter © 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2008.01.070
Fe–12Ni–6Mn steel during isothermal aging at 653 K. Later, Heo and Lee [5–7] reported a ductile–brittle–ductile transition in aged Fe–7Ni–8Mn and Fe–12Ni–6Mn steels which was well explained by successive manganese segregation and desegregation at grain boundaries during isothermal aging. Nevertheless, a few studies argued deleterious effect of manganese segregation on the intergranular failure of Fe–Ni–Mn steels. For instance, Suto and Murakami [8] identified that nickel and manganese concentrations at grain boundaries of an Fe–12Ni–6Mn steel never change in the ductile–brittle transition temperature and, consequently, criticized deleterious effect of manganese segregation in the intergranular failure! Alternatively, interaction of moving dislocations with disc-shaped precipitates at grain boundaries was suggested as a mechanism of grain boundary failure. Wayman and co-workers [9,10] did not find any evidences of second-phase particles or systematic correlation with manganese segregation at grain boundaries of embrittled Fe–20.8Ni–3.2Mn and Fe–20Ni–5Mn maraging steels. However, it has recently turned out that grain boundary precipitation is the main source of intergranular failure of Fe–Ni–Mn maraging steels. For instance, Mun et al. [11] reported precipitation of
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face-centered cubic (fcc) austenite particles at grain boundaries of an Fe–7Ni–8Mn steel at early stages of aging at 723 K for which decohesion of austenite–ferrite interface was proposed to augment grain boundary fracture. Further, Lee et al. [12] identified closely spaced precipitation of a face-centered tetragonal (fct) -NiMn intermetallic compound at grain boundaries of an Fe–10Ni–5Mn steel at early stages of aging at 753 K. Meanwhile, Wilson [13] argued that manganese segregates at grain boundaries as an initial step in the formation of grain boundary precipitates and acts as a major embrittling element in the early stages of aging. Therefore, the source of grain boundary embrittlement in Fe–Ni–Mn maraging steels has remained controversial yet. Hossein Nedjad studied an Fe–10Ni–7Mn maraging steel [14]. Intergranular embrittlement, grain boundary precipitation behavior and age-hardening precipitates of this steel during isothermal aging at 753 K have been reported already [15–18]. This paper is aimed to correlate the grain boundary embrittlement and precipitation reactions during isothermal aging treatment. 2. Experimental procedure An Fe–10.35Ni–6.88Mn–0.006C–0.007S–0.005P–0.005N– 0.003Al steel weighing 6 kg was prepared in a vacuum induction melting furnace under 10−2 mbar using electrolytic iron, electrolytic manganese and pure nickel shots. Bars weighing 200 g were cut from the ingot and remelted under argon gas in the water-cooled copper mold of a vacuum arc melting furnace. Remelted bars were encapsulated in quartz tubes under argon gas after evacuation to 10−5 mbar. Homogenizing treatment was performed at 1473 K for 173 ks followed by water quenching. Cold rolling to 85 pct reduction was carried out at room temperature followed by solution annealing treatment at 1223 K for 3.6 ks in a vacuum furnace, water quenching and subzero treatment at 77 K for 3.6 ks. Sheet-type tensile test pieces of 1 mm thickness, 2.5 mm width and 8.5 mm gage length were cut according to JIS Z2201 from a solution-annealed steel and aged for various times at 753 K in a neutralized salt bath. Tensile tests were carried out using a Shimadzu universal machine at a cross-head speed of 1 mm/min at room temperature. Fractography of broken tensile test pieces was performed by a scanning electron microscope. Auger electron spectroscopy studies of the solution-annealed and aged steels were carried out by a Physical Electronics1 PHI 680 scanning Auger microscope operating at a voltage of 10 kV and current of 10 nA. For transmission electron microscopy, disc-shaped specimens of diameter 3 mm and initial thickness 300 m were cut using an electro-discharge wire cutting machine, then manually polished to a thickness of ca. 30 m. For observation of deformed structure, rectangular specimens of dimension 0.3 mm × 1 mm × 3 mm were cut from a broken tensile test piece close to the fracture tip by an electro-discharge wire cutting machine, then manually
Fig. 1. Changes in the hardness, fracture stress and tensile elongation of the studied steel vs. isothermal aging time at 753 K.
polished to a thickness of ca. 30 m. Further thinning of thin foils was accomplished electrochemically in a solution of CrO3 (200 g), CH3 COOH (500 ml) and H2 O (40 ml) at 285 K using a TenuPol2 -3 instrument. Transmission electron microscopy was carried out with a PHILIPS3 CM200-FEG microscope operating at 200 kV. 3. Results Fig. 1 shows changes in the hardness, fracture stress and tensile elongation of the studied steel vs. isothermal aging time at 753 K. Hardness increases by increasing aging time up to a maximum of 585 Hv at 3.6 ks then decreases at later stages of aging. Fracture stress turns to increase at initial stages which, unexpectedly, decreases after short aging to a minimum at 0.36 ks. Then it increases at intermediate stages of aging and eventually decreases at later stages of aging. Tensile elongation decreases drastically at early stages of aging to zero, remains zero at intermediate stages and eventually increases with aging time at later stages of aging, resuming to a maximum of about 5 pct at prolonged aging times. Tensile test pieces aged at initial and intermediate stages were broken suddenly before yield point. Overaged specimens were also broken suddenly, but in a midway between yield strength and ultimate tensile strength. Fig. 2(a) shows a scanning electron micrograph of a broken solution-annealed tensile test piece indicating an intragranular dimpled ductile fracture. Scanning electron micrographs of specimens aged for 0.36 (underaged), 3.6 (peakaged) and 86.4 (overaged) ks are shown in Fig. 2(b), (c) and (d), respectively. Those broken tensile test pieces exhibit intergranular fracture with prevailing secondary cracks propagated along grain boundaries as indicated by dashed lines in Fig. 2(c). It is found out that the proportion of intergranular fracture increases qualitatively at early stages of aging and decreases after prolonged aging. High-magnification scanning electron micrographs illus2
1
Physical Electronics is a trademark of Physical Electronic Inc., Chanhassen, MN.
3
NJ.
TenuPol is a trademark of Struers A/S, Denmark. PHILIPS is a trademark of Philips Electronic Instruments Corp., Mahawh,
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Fig. 2. Scanning electron micrographs showing fracture surfaces of broken tensile test pieces; (a) solution-annealed; (b) underaged for 0.36 ks; (c) peakaged for 3.6 ks; (d) overaged for 86.4 ks. The solution-annealed and aged steels show ductile intragranular and brittle intergranular fracture modes, respectively. Dashed lines in (c) represent secondary grain boundary cracks.
trating enlargement of the break-opened grain boundaries at fracture surfaces of a solution-annealed specimen, broken intergranulary at 77 K, along with specimens aged for 0.36, 3.6 and 86.4 ks are shown in Fig. 3. Grain boundary fracture surface of the solution-annealed specimen in Fig. 3(a) seems microscopically smooth irrespective of rare striations. In the aged steels, Fig. 3(b), (c) and (d), grain boundary surfaces are microscopically rough in which grain boundary precipitates and striated features are found. An arrow in Fig. 4(d) denotes an striated feature which is proposed to result from decelerating microplastic overloading of facing matrices ahead of an advancing crack passing along grain boundary. However, such a striation could also be an array of precipitates at a lath boundary–grain boundary junction line, a lath boundary trace or an austenite lath precipitated at grain boundary, depending on the location and aging condition. It is found out that grain boundary precipitates become large and striations increase with increasing of aging time. Fig. 4(a) shows a high-magnification secondary electron micrograph of a break-opened grain boundary fracture surface of a specimen aged for 3.6 ks, which shows grain boundary precipitates settled in a facing matrix of a break-opened grain boundary. Auger electron spectroscopy of areas of about 1 m2 at grain boundary fracture surfaces of a solution-annealed specimen (the area shown by a rectangle in Fig. 3(a)) and a specimen aged for 3.6 ks (whole area of Fig. 4(a)) gave quantitative surface compositions of Fe–9.7Ni–8.6Mn and Fe–10.9Ni–15.5Mn, that is nickel and
manganese concentrations have increased in the aged steel which could be simply interpreted as nickel and manganese segregations at grain boundaries of the aged steel. Fig. 4(b) and (c) shows Auger electron maps corresponding to the grain boundary fracture surface shown in Fig. 4(a), illustrating nickel and manganese distributions at grain boundary, respectively. It shows clearly that grain boundary precipitates have been enriched by nickel and manganese atoms, but surrounding matrix have been depleted from those elements. Therefore, outstanding segregations of manganese and nickel atoms are unlikely at grain boundaries of the aged steel. Further, the grain boundary enrichment outlined by Auger electron spectroscopy of the aged steel is attributed to nickel and manganese atoms accumulated at grain boundary precipitates, not to any mono/multilayer atomic segregations where grain boundary precipitates could drain preliminary-segregated manganese and nickel atoms out of grain boundaries. Fig. 5(a) and (b) shows dark-field transmission electron micrographs obtained using spots corresponding to the {1 1 1}planes of the fct -NiMn intermetallic phase in zone axes close to the 1 0 0 directions of bcc iron matrix in steels aged for 0.36 and 86.4 ks, respectively. In the early stages of aging, nanometer-scaled particulates can be deduced which have grown to disc-shaped precipitates at later stages of aging. Fig. 6(a) shows a bright-field transmission electron micrograph of a grain boundary in a specimen aged for 3.6 ks in which coarse grain
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Fig. 3. High-magnification scanning electron micrographs showing enlargement of the break-opened grain boundaries; (a) solution-annealed and broken at 77 K; (b) underaged for 0.36 ks; (c) peakaged for 3.6 ks; (d) overaged for 86.4 ks. Fracture surface of the solution-annealed steel is microscopically smooth but, in the aged steels grain boundary precipitates and striations (arrowed in 3d, for example) are identified.
boundary precipitates in association with precipitate-free zone is demonstrated. A bright-field transmission electron micrograph of a grain boundary in a broken tensile test piece of the specimen aged for 3.6 ks is shown in Fig. 6(b), illustrating strain configuration during tensile deformation, raised at matrix and a grain boundary with coarse precipitates and precipitate-free zone. Within the upper grain, inhomogeneous deformation band is observed (see a set of wavy lines) where at the end of the inhomogeneously deformed region, microscopic strain localization has been created at grain boundary (inside rectangle). An enlargement of the strain localization area shown by the rectangle in Fig. 6(b), is shown in Fig. 6(c) wherein bands of lamellar fringes with about 20 nm width (denoted by symbols B1 and B2) are observed, lying in a direction inclined at about thirty degrees to the grain boundary line. Those bands are proposed to be stacking faults, realizing inhomogeneous deformation of bcc iron matrix in terms of the so-called planar slip. The planar slip bands have propagated to relatively long ranges at bcc iron matrix, evidencing that matrix precipitates have been sheared by moving partial dislocations creating the stacking faults (see inset, showing an enlargement of the band B1). Fig. 7(a) and (b) shows bright-field transmission electron micrographs evidencing strain localization at precipitate-free zone resulted from the matrix slip in a broken tensile test piece of the specimen aged for 3.6 ks. Fig. 7(a) demonstrates higher magnification of the rectangled area of Fig. 6(b), wherein strain localization at
precipitate-free zone is indicated by an arrow and the arrow tips show crack initiation at grain boundary. Dislocation tangles created at matrix adjacent to a grain boundary in Fig. 7(b) shows that remarked strain localization occurs as a result of inhomogeneous matrix deformation in the loaded tensile test piece with capabilities to fascinate crack initiation at grain boundary. 4. Discussion It is found out that the Fe–10Ni–7Mn maraging steel suffers from grain boundary fracture after aging treatment. The intergranular brittleness causes drastic decreasing of both fracture stress and tensile ductility such that in the underaged and peakaged steels, fracture occurs catastrophically with zero tensile elongation before yield point. Fractographic examination of fracture surfaces revealed coarse grain boundary precipitates at break-opened grain boundaries along with striated features and holes of pulled out grain boundary precipitates. On the other hand, transmission electron microscopy of aged specimens demonstrated formation of coarse precipitates associated with precipitate-free zones at grain boundaries, addressing the weakening of grain boundaries by coarsening reaction. A detailed description of the grain boundary precipitation reaction of the present steel has been published elsewhere [16]. However, after tensile deformation of a peakaged specimen, inhomogeneous deformation in terms of the planar slip bands was found at matrix
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Fig. 4. (a) High-magnification secondary electron image of a break-opened grain boundary showing precipitates in a steel aged for 3.6 ks; (b) Auger electron map of nickel; and (c) manganese corresponding to the scanning electron micrograph.
where at the end of the planar slip band, localized straining of precipitate-free zone and initiation of grain boundary microcracks were illustrated. Upon further straining, those microcracks were assumed to propagate along grain boundaries, breakopening the soft precipitate-free zone which is low strength bcc iron (ferrite) phase with capability to act as a microscopically ductile phase in front of the advancing cracks. Consequently, fracture could proceed by virtue of microplastic rupture in the soft precipitate-free zone and, therefore, a ductile grain boundary fracture mechanism could be suggested for intergranular failure of the present steel. Grain boundary ductile fracture of precipitation hardening alloys has been recognized for a long time [19]. The phenomenon is quite different from the brittle fracture arising from segregation of harmful impurities at grain boundaries which usually gives microscopically smooth fracture surfaces [20]. Although manganese segregation has been suggested as a source of grain boundary fracture in Fe–Ni–Mn steels already, but operation of the segregation-induced brittleness in the present steel is unlikely due to the presence of grain boundary fct -NiMn precipitates which drain nickel and manganese atoms out of grain
boundaries. However, three modes have been introduced for grain boundary ductile fracture in the precipitation hardening alloys [21]; (i) microvoid nucleation at large grain boundary precipitates, not associated with precipitate-free zone, due to stress–strain incompatibilities in the presence of large nondeformable grain boundary precipitates, (ii) strain localization in the soft precipitate-free zone associated with grain boundary precipitation and (iii) shearing of matrix precipitates by moving dislocations, giving rise to the inhomogeneous planar slip bands at grain interior that may terminate at grain boundaries and, subsequently, apply large stress concentrations to augment grain boundary fracture. In this case, for break-opening of a grain boundary to result from planar slip stress concentration, the grain boundary must become weakened otherwise strain concentration can be relived by shear in the adjacent grain without grain boundary fracture. Operation of the planar slip mode of fracture in the peakaged specimen of the present steel was demonstrated. Generally, planar slip is more likely in metallic solid solution of low stacking-fault energy like copper–aluminum [22] and copper–zinc [23] alloys. Nevertheless, operation of planar slip
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Fig. 5. (a) Dark-field transmission electron micrograph obtained using the {1 1 1}fct -NiMn phase showing fine precipitates in a steel (a) underaged for 0.36 ks; (b) overaged for 86.4 ks.
has also been reported in metals with high stacking-fault energy, e.g. aluminium–lithium alloys [21,24]. Further, age-hardening alloys with shearable second-phase precipitates at matrix also exhibit planar slip in which cutting of precipitates by moving dislocations reduces their effective size, resulting in weakening of the obstacles and deformation localization [25]. Often, there are other factors which stimulate strain localization in the form of planar slip bands, e.g. (i) internal ordering or microtwinning of the second-phase particles in age-hardening alloys [26,27], (ii) short range ordering or clustering of matrix [26], and (iii) cyclic mode of deformation [28]. In the Fe–Ni–Mn steel, formation of ordered (B2, CsCl type) bcc -NiMn zones due to a miscibility gap has been suggested for early stages of aging [29]. Formation of nanometer-sized fct -NiMn precipitates has been identified in an underaged Fe–10Ni–Mn steel which is proposed to form by conversion of the preliminarily formed bcc -NiMn zones in an underaging state before peak hardness. Further, those fct -NiMn precipitates have been found to be ordered (L10 , CuAl type) and internally microtwinned on {1 1 1}fct planes in response to the transformation or coherency strain accommodation [18]. Those nanometer-sized bcc -NiMn zones as well as the nanometer-sized fct -NiMn precipitates are likely to be sheared by dislocations in the underaged and peakaged steels which may realize operation of the planar slip in the present bcc metal. Further, as a perquisite for grain boundary ductile fracture of the present steel to result from planar slip, weaken-
Fig. 6. Bright-field transmission electron micrograph showing (a) coarse grain boundary precipitates () in association with precipitate-free zone (PFZ) at a grain boundary (PAGB) of a specimen peakaged for 3.6 ks; (b) after tensile deformation showing inhomogeneous deformation band (parallel white lines) at upper grain along with strain localization at precipitate-free zone in the vicinity of a grain boundary precipitates (see inside rectangle); (c) higher magnification view of the rectangle area showing planar slip bands (B1 and B2). See inset for interference fringes of the band B1as an indicative of stacking faults in bcc iron matrix.
ing of grain boundaries by discontinuous coarsening reaction is pointed out. Inasmuch as the discontinuous coarsening aggregate at grain boundary has lower strength than matrices of the facing grains, propagating cracks are more likely to break-open grain boundary than to shear facing matrix by a grain boundary offset. In the underaged steel, those grain boundary precipitates are in the form of platelets located in an inclined planar array with respect to grain boundary plane (see Fig. 8 of Ref. [16]). Such planar arrays of precipitates are proposed to raise stress concentration at the edges of grain boundary precipitates as a result of which critical stress or strain for crack initiation is reduced in the underaged steel. By further aging, matrix precipitate overage by continuous coarsening and their resistance against shearing increases. It has been well established that dislocations bypass rather than shear precipitates beyond the peak in yield stress and consequently, planar slip mode of fracture no longer operates
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Fig. 7. Bright-field transmission electron micrographs evidencing strain localization at precipitate-free zone. (a) Enlargement of rectangled area in Fig. 6(b) wherein arrow and the arrow tips denote strain localization and microcracks initiation at grain boundary, respectively; (b) arrows indicate dislocation tangles and strain contrast in the vicinity of grain boundary.
at matrix to be substituted by conventional multi-system slip [21,25]. Consequently, grain boundary fracture at and beyond peak aging will be then determined by strain localization in the soft precipitate-free zone. At the same time, overaging of matrix precipitates decreases yield strength of matrix as a result of which matrix deformation incorporates into overall deformation process which is likely to increase tensile ductility and the proportion of intragranular fracture. Further, grain boundary precipitates coarsen during isothermal aging and, hence, the area fraction of grain boundary precipitates decreases by increasing of aging time. It has been well known that the proportion of intragranular fracture and fracture stress increase with decreasing the area fraction of grain boundary precipitates [30]. Furthermore, the preliminary planar array of plate-shaped grain boundary precipitates is modified to spherical morphologies at later stages of aging by virtue of a grain boundary diffusionassisted coarsening reaction, which is expected to decrease stress concentration and, subsequently, increase critical stress or strain for microcrack initiation at precipitate-free zone in the vicinity of grain boundary precipitates. Simultaneously, precipitate-free zone increases in width by increasing aging time. Increasing of precipitate-free zone width leads to decreasing of grain boundary fracture stress if the width of zones exceeds considerably the interparticle spacing of matrix precipitates [31]. Tensile ductil-
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ity also increases slightly as the precipitate-free zone increases in width [30]. In the present steel, width of precipitate-free zone is considerably larger than the spacing of matrix precipitates in all aging times studied. Therefore, presence of precipitate-free zone at grain boundaries of the present steel should be detrimental on the grain boundary fracture of the present steel, tough it may incorporate marginally into tensile ductility at prolonged aging. During further aging beyond peak hardness, overaging of matrix precipitates, coarsening of grain boundary precipitates associated with morphological modification and increasing of precipitate-free zone width occur concurrently. Diffusional transformation into thermodynamically stable fcc austenite also takes place mainly at grain and lath boundaries [17]. A combination of the aforementioned evolutions is expected to increase tensile ductility by increasing of aging time. Nevertheless, tensile ductility still remains zero at intermediate stages up to a few hours. Therefore, it seems that activation of matrix planar slip and coarse grain boundary precipitates in association with precipitate-free zone offset the beneficial microstructural evolutions to improve ductility. After so prolonged aging for larger than 1 day, austenite becomes predominant phase which incorporates to deformation effectively and, consequently, decreases fracture stress and raises ductility by arresting the propagating cracks. Lee et al. [12] attributed the grain boundary fracture of an Fe–10Ni–5Mn maraging steel to closely spaced grain boundary precipitates for which stress–strain incompatibilities at grain boundaries due to intermetallic precipitates could be assumed as an operative mechanism. However, in the present steel grain boundary precipitation was recognized in association with precipitate-free zone. Furthermore, synergistically detrimental effects of inhomogeneous matrix slip and grain boundary weakening were proposed to realize catastrophic grain boundary failure in the early stages of aging. However, more studies are necessary to clarify matrix deformation characteristics and to prohibit grain boundary precipitation reactions of Fe–Ni–Mn steels to suppress their intergranular failure. 5. Conclusions Evolution of the characteristics of grain boundary fracture in correlation with precipitation reactions during isothermal aging at 753 K was studied in an Fe–10Ni–7Mn maraging steel. Main conclusions drawn were as follows: 1. The Fe–10Ni–7Mn maraging steel shows substantial hardening but suffers from premature intergranular failure after aging. 2. Fracture stress turns to increase at initial stages of aging, which unexpectedly decreases after short aging. After a minimum at short aging, fracture stress increases at intermediate stages of aging and eventually decreases at later stages of aging. Tensile elongation decreases drastically at early stages of aging to zero which resumes to a few pct after prolonged aging. 3. A microscopically ductile grain boundary fracture was identified in the aged steel which seems macroscopically brittle
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from the premature onset of fracture and zero tensile elongation. 4. Heterogeneous grain boundary precipitation occurs during isothermal aging in association with solute-depleted precipitate-free zone, realizing a network of weakened material along grain boundaries in the bulk. 5. Inhomogeneous matrix deformation in the form of planar slip bands was found to raise strain localization at precipitate-free zone which reduces critical strain for crack initiation at grain boundary. 6. After prolonged aging, overaging of matrix precipitates, morphological modification of grain boundaries, growth of precipitate-free zone in width increase tensile ductility in association with isothermally transformed austenite particles. References [1] M. Tanaka, T. Suzuki, M. Yodogawa, Bull. Tokyo Inst. Tech. 82 (1967) 79–90. [2] M. Tanaka, J. Yamamoto, Toward Improved Ductility and Toughness, Climax Mo. Dev. Co, Japan, 1971, pp. 195–206. [3] D.R. Squires, E.A. Wilson, Metall. Trans. 3 (1972) 575–581. [4] H.C. Feng, E.A. Wilson, C.J. McMahon Jr., 3rd Int. Conf. on the Strength of Metals and Steels, Cambridge, England, 1973, pp. 129–133. [5] N.H. Heo, H.C. Lee, Metall. Trans. 27A (1996) 1015–1020. [6] N.H. Heo, Acta Mater. 44 (1996) 3015–3023. [7] N.H. Heo, Scripta Mater. 34 (1996) 1517–1522. [8] H. Suto, T. Murakami, Trans. JIM 20 (1979) 365–370.
[9] S.J. Kim, C.M. Wayman, Mater. Sci. Eng. 136A (1991) 121–129. [10] L.T. Shiang, C.M. Wayman, Metallography 21 (1988) 399–423. [11] S.H. Mun, M. Watanabe, X. Li, K.H. OH, D.B. Williams, H.C. Lee, Metall. Mater. Trans. 33A (2002) 1057–1067. [12] H.C. Lee, S.H. Mun, D. McKenzie, Metall. Mater. Trans. 34A (2003) 2421–2428. [13] E.A. Wilson, Metall. Mater. Trans. 35A (2004) 352–355. [14] S. Hossein Nedjad, Ph.D. Thesis, University of Tehran, Tehran, 2005. [15] S. Hossein Nedjad, M. Nili Ahmadabadi, R. Mahmudi, H. Farhangi, Mater. Sci. Eng. A 378 (2004) 314–318. [16] S. Hossein Nedjad, M. Nili Ahmadabadi, T. Furuhara, Metall. Mater. Trans. A 39A (2008) 19–27. [17] S. Hossein Nedjad, M. Nili Ahmadabadi, R. Mahmudi, T. Furuhara, T. Maki, Mater. Sci. Eng. 438–440A (2006) 288–291. [18] S. Hossein Nedjad, M. Nili Ahmadabadi, T. Furuhara, T. Maki, Phys. Status Solidi (a) 203 (2006) 2229–2235. [19] A.H. Geisler, in: R. Smoluchowski (Ed.), Phase Transformations in Solids, Wiley, New York, 1951. [20] E. Hondros, M.P. Seah, Int. Metall. Rev. 222 (1977) 262–301. [21] A.K. Vasudevan, R.D. Doherty, Acta Metall. 35 (1987) 1193–1219. [22] B.D. Yan, A.S. Cheng, L. Buchinger, S. Stanzl, C. Laird, Mater. Sci. Eng. 80 (1986) 129–142. [23] P. Lukas, L. Kunz, J. Krejci, Mater. Sci. Eng. 158A (1992) 177–183. [24] A.A. Csontos, E.A. Starke, Int. J. Plast. 21 (2005) 1097–1118. [25] J.W. Martin, Precipitation Hardening, second ed., Butterworth-Heinemann, Great Britain, 1998. [26] V. Gerold, H.P. Karnthaler, Acta Metall. 37 (1989) 2177–2183. [27] H.F. Merrick, Metall. Trans. A 7A (1976) 505–514. [28] L. Xiao, D.L. Chen, M.C. Chaturvedi, Scripta Mater. 52 (2005) 603–607. [29] T. Suzuki, Trans. JIM 11 (1970) 257–263. [30] P.T. Unwin, G.C. Smith, J. Inst. Met. 97 (1969) 299–310. [31] T. Krol, D. Baither, E. Nembach, Scripta Mater. 48 (2003) 1189–1194.
Correlation between the intergranular brittleness and precipitation reactions during isothermal aging of an Fe–Ni–Mn maraging steel S. Hossein Nedjad a,∗ , M. Nili Ahmadabadi b , T. Furuhara c a b
Faculty of Materials Engineering, Sahand University of Technology, P.O. Box 51335-1996, Tabriz, Iran School of Metallurgy and Materials Engineering, University of Tehran, P.O. Box 14395-731, Tehran, Iran c Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan Received 27 August 2007; received in revised form 25 December 2007; accepted 24 January 2008
Abstract Evolution of the intergranular brittleness of an Fe–10Ni–7Mn (weight pct) maraging steel was correlated with precipitation reactions during isothermal aging at 753 K. Intergranular brittleness of the Fe–Ni–Mn steel raises after aging treatment which occurs catastrophically at zero tensile elongation in the underaged and peakaged steels. The intergranular failure is attributed to grain boundary weakening due to the formation of coarse grain boundary precipitates associated with solute-depleted precipitate-free zones during isothermal aging. Further, evidences of planar slip bands were found within the grains of a peakaged specimen loaded by tensile deformation. Those inhomogeneously deformed bands were identified to apply large strain localization in the soft precipitate-free zones at grain boundaries which is assumed to fascinate microcracks initiation at negligible macroscopic strains in the underaged and peakaged steels. During further aging, concurrent reactions including (i) overaging of matrix precipitates, (ii) spheroidization of grain boundary precipitates, (iii) growth of precipitate-free zone in width and (iv) diffusional transformation to austenite take place which increase tensile ductility after prolonged aging. © 2008 Elsevier B.V. All rights reserved. Keywords: Maraging steel; Intergranular brittleness; Grain boundary precipitation; Precipitate-free zone; Planar slip
1. Introduction Iron–nickel–manganese martensitic steels show substantial age hardening but suffer from poor ductility after aging [1,2]. The Fe–Ni–Mn maraging steels exhibit intragranular dimpled ductile fracture in the solution-annealed condition which turns, by short isothermal aging, to premature intergranular brittle fracture, passing along prior austenite grain boundaries (PAGBs) [3]. Hereafter, the PAGBs will be refereed to as grain boundaries. Analogous to the temper-embrittlement of low alloy steels, Squires and Wilson [3] first suggested that intergranular failure of an Fe–12Ni–6Mn (weight pct) maraging steel arises from segregation of manganese at grain boundaries during isothermal aging at 573–773 K. Hereafter, all chemical compositions will be given in weight pct. Then Feng et al. [4] showed by Auger electron spectroscopy (AES) that nickel and manganese segregate at grain boundaries of an
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Fe–12Ni–6Mn steel during isothermal aging at 653 K. Later, Heo and Lee [5–7] reported a ductile–brittle–ductile transition in aged Fe–7Ni–8Mn and Fe–12Ni–6Mn steels which was well explained by successive manganese segregation and desegregation at grain boundaries during isothermal aging. Nevertheless, a few studies argued deleterious effect of manganese segregation on the intergranular failure of Fe–Ni–Mn steels. For instance, Suto and Murakami [8] identified that nickel and manganese concentrations at grain boundaries of an Fe–12Ni–6Mn steel never change in the ductile–brittle transition temperature and, consequently, criticized deleterious effect of manganese segregation in the intergranular failure! Alternatively, interaction of moving dislocations with disc-shaped precipitates at grain boundaries was suggested as a mechanism of grain boundary failure. Wayman and co-workers [9,10] did not find any evidences of second-phase particles or systematic correlation with manganese segregation at grain boundaries of embrittled Fe–20.8Ni–3.2Mn and Fe–20Ni–5Mn maraging steels. However, it has recently turned out that grain boundary precipitation is the main source of intergranular failure of Fe–Ni–Mn maraging steels. For instance, Mun et al. [11] reported precipitation of
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face-centered cubic (fcc) austenite particles at grain boundaries of an Fe–7Ni–8Mn steel at early stages of aging at 723 K for which decohesion of austenite–ferrite interface was proposed to augment grain boundary fracture. Further, Lee et al. [12] identified closely spaced precipitation of a face-centered tetragonal (fct) -NiMn intermetallic compound at grain boundaries of an Fe–10Ni–5Mn steel at early stages of aging at 753 K. Meanwhile, Wilson [13] argued that manganese segregates at grain boundaries as an initial step in the formation of grain boundary precipitates and acts as a major embrittling element in the early stages of aging. Therefore, the source of grain boundary embrittlement in Fe–Ni–Mn maraging steels has remained controversial yet. Hossein Nedjad studied an Fe–10Ni–7Mn maraging steel [14]. Intergranular embrittlement, grain boundary precipitation behavior and age-hardening precipitates of this steel during isothermal aging at 753 K have been reported already [15–18]. This paper is aimed to correlate the grain boundary embrittlement and precipitation reactions during isothermal aging treatment. 2. Experimental procedure An Fe–10.35Ni–6.88Mn–0.006C–0.007S–0.005P–0.005N– 0.003Al steel weighing 6 kg was prepared in a vacuum induction melting furnace under 10−2 mbar using electrolytic iron, electrolytic manganese and pure nickel shots. Bars weighing 200 g were cut from the ingot and remelted under argon gas in the water-cooled copper mold of a vacuum arc melting furnace. Remelted bars were encapsulated in quartz tubes under argon gas after evacuation to 10−5 mbar. Homogenizing treatment was performed at 1473 K for 173 ks followed by water quenching. Cold rolling to 85 pct reduction was carried out at room temperature followed by solution annealing treatment at 1223 K for 3.6 ks in a vacuum furnace, water quenching and subzero treatment at 77 K for 3.6 ks. Sheet-type tensile test pieces of 1 mm thickness, 2.5 mm width and 8.5 mm gage length were cut according to JIS Z2201 from a solution-annealed steel and aged for various times at 753 K in a neutralized salt bath. Tensile tests were carried out using a Shimadzu universal machine at a cross-head speed of 1 mm/min at room temperature. Fractography of broken tensile test pieces was performed by a scanning electron microscope. Auger electron spectroscopy studies of the solution-annealed and aged steels were carried out by a Physical Electronics1 PHI 680 scanning Auger microscope operating at a voltage of 10 kV and current of 10 nA. For transmission electron microscopy, disc-shaped specimens of diameter 3 mm and initial thickness 300 m were cut using an electro-discharge wire cutting machine, then manually polished to a thickness of ca. 30 m. For observation of deformed structure, rectangular specimens of dimension 0.3 mm × 1 mm × 3 mm were cut from a broken tensile test piece close to the fracture tip by an electro-discharge wire cutting machine, then manually
Fig. 1. Changes in the hardness, fracture stress and tensile elongation of the studied steel vs. isothermal aging time at 753 K.
polished to a thickness of ca. 30 m. Further thinning of thin foils was accomplished electrochemically in a solution of CrO3 (200 g), CH3 COOH (500 ml) and H2 O (40 ml) at 285 K using a TenuPol2 -3 instrument. Transmission electron microscopy was carried out with a PHILIPS3 CM200-FEG microscope operating at 200 kV. 3. Results Fig. 1 shows changes in the hardness, fracture stress and tensile elongation of the studied steel vs. isothermal aging time at 753 K. Hardness increases by increasing aging time up to a maximum of 585 Hv at 3.6 ks then decreases at later stages of aging. Fracture stress turns to increase at initial stages which, unexpectedly, decreases after short aging to a minimum at 0.36 ks. Then it increases at intermediate stages of aging and eventually decreases at later stages of aging. Tensile elongation decreases drastically at early stages of aging to zero, remains zero at intermediate stages and eventually increases with aging time at later stages of aging, resuming to a maximum of about 5 pct at prolonged aging times. Tensile test pieces aged at initial and intermediate stages were broken suddenly before yield point. Overaged specimens were also broken suddenly, but in a midway between yield strength and ultimate tensile strength. Fig. 2(a) shows a scanning electron micrograph of a broken solution-annealed tensile test piece indicating an intragranular dimpled ductile fracture. Scanning electron micrographs of specimens aged for 0.36 (underaged), 3.6 (peakaged) and 86.4 (overaged) ks are shown in Fig. 2(b), (c) and (d), respectively. Those broken tensile test pieces exhibit intergranular fracture with prevailing secondary cracks propagated along grain boundaries as indicated by dashed lines in Fig. 2(c). It is found out that the proportion of intergranular fracture increases qualitatively at early stages of aging and decreases after prolonged aging. High-magnification scanning electron micrographs illus2
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Physical Electronics is a trademark of Physical Electronic Inc., Chanhassen, MN.
3
NJ.
TenuPol is a trademark of Struers A/S, Denmark. PHILIPS is a trademark of Philips Electronic Instruments Corp., Mahawh,
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Fig. 2. Scanning electron micrographs showing fracture surfaces of broken tensile test pieces; (a) solution-annealed; (b) underaged for 0.36 ks; (c) peakaged for 3.6 ks; (d) overaged for 86.4 ks. The solution-annealed and aged steels show ductile intragranular and brittle intergranular fracture modes, respectively. Dashed lines in (c) represent secondary grain boundary cracks.
trating enlargement of the break-opened grain boundaries at fracture surfaces of a solution-annealed specimen, broken intergranulary at 77 K, along with specimens aged for 0.36, 3.6 and 86.4 ks are shown in Fig. 3. Grain boundary fracture surface of the solution-annealed specimen in Fig. 3(a) seems microscopically smooth irrespective of rare striations. In the aged steels, Fig. 3(b), (c) and (d), grain boundary surfaces are microscopically rough in which grain boundary precipitates and striated features are found. An arrow in Fig. 4(d) denotes an striated feature which is proposed to result from decelerating microplastic overloading of facing matrices ahead of an advancing crack passing along grain boundary. However, such a striation could also be an array of precipitates at a lath boundary–grain boundary junction line, a lath boundary trace or an austenite lath precipitated at grain boundary, depending on the location and aging condition. It is found out that grain boundary precipitates become large and striations increase with increasing of aging time. Fig. 4(a) shows a high-magnification secondary electron micrograph of a break-opened grain boundary fracture surface of a specimen aged for 3.6 ks, which shows grain boundary precipitates settled in a facing matrix of a break-opened grain boundary. Auger electron spectroscopy of areas of about 1 m2 at grain boundary fracture surfaces of a solution-annealed specimen (the area shown by a rectangle in Fig. 3(a)) and a specimen aged for 3.6 ks (whole area of Fig. 4(a)) gave quantitative surface compositions of Fe–9.7Ni–8.6Mn and Fe–10.9Ni–15.5Mn, that is nickel and
manganese concentrations have increased in the aged steel which could be simply interpreted as nickel and manganese segregations at grain boundaries of the aged steel. Fig. 4(b) and (c) shows Auger electron maps corresponding to the grain boundary fracture surface shown in Fig. 4(a), illustrating nickel and manganese distributions at grain boundary, respectively. It shows clearly that grain boundary precipitates have been enriched by nickel and manganese atoms, but surrounding matrix have been depleted from those elements. Therefore, outstanding segregations of manganese and nickel atoms are unlikely at grain boundaries of the aged steel. Further, the grain boundary enrichment outlined by Auger electron spectroscopy of the aged steel is attributed to nickel and manganese atoms accumulated at grain boundary precipitates, not to any mono/multilayer atomic segregations where grain boundary precipitates could drain preliminary-segregated manganese and nickel atoms out of grain boundaries. Fig. 5(a) and (b) shows dark-field transmission electron micrographs obtained using spots corresponding to the {1 1 1}planes of the fct -NiMn intermetallic phase in zone axes close to the 1 0 0 directions of bcc iron matrix in steels aged for 0.36 and 86.4 ks, respectively. In the early stages of aging, nanometer-scaled particulates can be deduced which have grown to disc-shaped precipitates at later stages of aging. Fig. 6(a) shows a bright-field transmission electron micrograph of a grain boundary in a specimen aged for 3.6 ks in which coarse grain
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Fig. 3. High-magnification scanning electron micrographs showing enlargement of the break-opened grain boundaries; (a) solution-annealed and broken at 77 K; (b) underaged for 0.36 ks; (c) peakaged for 3.6 ks; (d) overaged for 86.4 ks. Fracture surface of the solution-annealed steel is microscopically smooth but, in the aged steels grain boundary precipitates and striations (arrowed in 3d, for example) are identified.
boundary precipitates in association with precipitate-free zone is demonstrated. A bright-field transmission electron micrograph of a grain boundary in a broken tensile test piece of the specimen aged for 3.6 ks is shown in Fig. 6(b), illustrating strain configuration during tensile deformation, raised at matrix and a grain boundary with coarse precipitates and precipitate-free zone. Within the upper grain, inhomogeneous deformation band is observed (see a set of wavy lines) where at the end of the inhomogeneously deformed region, microscopic strain localization has been created at grain boundary (inside rectangle). An enlargement of the strain localization area shown by the rectangle in Fig. 6(b), is shown in Fig. 6(c) wherein bands of lamellar fringes with about 20 nm width (denoted by symbols B1 and B2) are observed, lying in a direction inclined at about thirty degrees to the grain boundary line. Those bands are proposed to be stacking faults, realizing inhomogeneous deformation of bcc iron matrix in terms of the so-called planar slip. The planar slip bands have propagated to relatively long ranges at bcc iron matrix, evidencing that matrix precipitates have been sheared by moving partial dislocations creating the stacking faults (see inset, showing an enlargement of the band B1). Fig. 7(a) and (b) shows bright-field transmission electron micrographs evidencing strain localization at precipitate-free zone resulted from the matrix slip in a broken tensile test piece of the specimen aged for 3.6 ks. Fig. 7(a) demonstrates higher magnification of the rectangled area of Fig. 6(b), wherein strain localization at
precipitate-free zone is indicated by an arrow and the arrow tips show crack initiation at grain boundary. Dislocation tangles created at matrix adjacent to a grain boundary in Fig. 7(b) shows that remarked strain localization occurs as a result of inhomogeneous matrix deformation in the loaded tensile test piece with capabilities to fascinate crack initiation at grain boundary. 4. Discussion It is found out that the Fe–10Ni–7Mn maraging steel suffers from grain boundary fracture after aging treatment. The intergranular brittleness causes drastic decreasing of both fracture stress and tensile ductility such that in the underaged and peakaged steels, fracture occurs catastrophically with zero tensile elongation before yield point. Fractographic examination of fracture surfaces revealed coarse grain boundary precipitates at break-opened grain boundaries along with striated features and holes of pulled out grain boundary precipitates. On the other hand, transmission electron microscopy of aged specimens demonstrated formation of coarse precipitates associated with precipitate-free zones at grain boundaries, addressing the weakening of grain boundaries by coarsening reaction. A detailed description of the grain boundary precipitation reaction of the present steel has been published elsewhere [16]. However, after tensile deformation of a peakaged specimen, inhomogeneous deformation in terms of the planar slip bands was found at matrix
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Fig. 4. (a) High-magnification secondary electron image of a break-opened grain boundary showing precipitates in a steel aged for 3.6 ks; (b) Auger electron map of nickel; and (c) manganese corresponding to the scanning electron micrograph.
where at the end of the planar slip band, localized straining of precipitate-free zone and initiation of grain boundary microcracks were illustrated. Upon further straining, those microcracks were assumed to propagate along grain boundaries, breakopening the soft precipitate-free zone which is low strength bcc iron (ferrite) phase with capability to act as a microscopically ductile phase in front of the advancing cracks. Consequently, fracture could proceed by virtue of microplastic rupture in the soft precipitate-free zone and, therefore, a ductile grain boundary fracture mechanism could be suggested for intergranular failure of the present steel. Grain boundary ductile fracture of precipitation hardening alloys has been recognized for a long time [19]. The phenomenon is quite different from the brittle fracture arising from segregation of harmful impurities at grain boundaries which usually gives microscopically smooth fracture surfaces [20]. Although manganese segregation has been suggested as a source of grain boundary fracture in Fe–Ni–Mn steels already, but operation of the segregation-induced brittleness in the present steel is unlikely due to the presence of grain boundary fct -NiMn precipitates which drain nickel and manganese atoms out of grain
boundaries. However, three modes have been introduced for grain boundary ductile fracture in the precipitation hardening alloys [21]; (i) microvoid nucleation at large grain boundary precipitates, not associated with precipitate-free zone, due to stress–strain incompatibilities in the presence of large nondeformable grain boundary precipitates, (ii) strain localization in the soft precipitate-free zone associated with grain boundary precipitation and (iii) shearing of matrix precipitates by moving dislocations, giving rise to the inhomogeneous planar slip bands at grain interior that may terminate at grain boundaries and, subsequently, apply large stress concentrations to augment grain boundary fracture. In this case, for break-opening of a grain boundary to result from planar slip stress concentration, the grain boundary must become weakened otherwise strain concentration can be relived by shear in the adjacent grain without grain boundary fracture. Operation of the planar slip mode of fracture in the peakaged specimen of the present steel was demonstrated. Generally, planar slip is more likely in metallic solid solution of low stacking-fault energy like copper–aluminum [22] and copper–zinc [23] alloys. Nevertheless, operation of planar slip
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Fig. 5. (a) Dark-field transmission electron micrograph obtained using the {1 1 1}fct -NiMn phase showing fine precipitates in a steel (a) underaged for 0.36 ks; (b) overaged for 86.4 ks.
has also been reported in metals with high stacking-fault energy, e.g. aluminium–lithium alloys [21,24]. Further, age-hardening alloys with shearable second-phase precipitates at matrix also exhibit planar slip in which cutting of precipitates by moving dislocations reduces their effective size, resulting in weakening of the obstacles and deformation localization [25]. Often, there are other factors which stimulate strain localization in the form of planar slip bands, e.g. (i) internal ordering or microtwinning of the second-phase particles in age-hardening alloys [26,27], (ii) short range ordering or clustering of matrix [26], and (iii) cyclic mode of deformation [28]. In the Fe–Ni–Mn steel, formation of ordered (B2, CsCl type) bcc -NiMn zones due to a miscibility gap has been suggested for early stages of aging [29]. Formation of nanometer-sized fct -NiMn precipitates has been identified in an underaged Fe–10Ni–Mn steel which is proposed to form by conversion of the preliminarily formed bcc -NiMn zones in an underaging state before peak hardness. Further, those fct -NiMn precipitates have been found to be ordered (L10 , CuAl type) and internally microtwinned on {1 1 1}fct planes in response to the transformation or coherency strain accommodation [18]. Those nanometer-sized bcc -NiMn zones as well as the nanometer-sized fct -NiMn precipitates are likely to be sheared by dislocations in the underaged and peakaged steels which may realize operation of the planar slip in the present bcc metal. Further, as a perquisite for grain boundary ductile fracture of the present steel to result from planar slip, weaken-
Fig. 6. Bright-field transmission electron micrograph showing (a) coarse grain boundary precipitates () in association with precipitate-free zone (PFZ) at a grain boundary (PAGB) of a specimen peakaged for 3.6 ks; (b) after tensile deformation showing inhomogeneous deformation band (parallel white lines) at upper grain along with strain localization at precipitate-free zone in the vicinity of a grain boundary precipitates (see inside rectangle); (c) higher magnification view of the rectangle area showing planar slip bands (B1 and B2). See inset for interference fringes of the band B1as an indicative of stacking faults in bcc iron matrix.
ing of grain boundaries by discontinuous coarsening reaction is pointed out. Inasmuch as the discontinuous coarsening aggregate at grain boundary has lower strength than matrices of the facing grains, propagating cracks are more likely to break-open grain boundary than to shear facing matrix by a grain boundary offset. In the underaged steel, those grain boundary precipitates are in the form of platelets located in an inclined planar array with respect to grain boundary plane (see Fig. 8 of Ref. [16]). Such planar arrays of precipitates are proposed to raise stress concentration at the edges of grain boundary precipitates as a result of which critical stress or strain for crack initiation is reduced in the underaged steel. By further aging, matrix precipitate overage by continuous coarsening and their resistance against shearing increases. It has been well established that dislocations bypass rather than shear precipitates beyond the peak in yield stress and consequently, planar slip mode of fracture no longer operates
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Fig. 7. Bright-field transmission electron micrographs evidencing strain localization at precipitate-free zone. (a) Enlargement of rectangled area in Fig. 6(b) wherein arrow and the arrow tips denote strain localization and microcracks initiation at grain boundary, respectively; (b) arrows indicate dislocation tangles and strain contrast in the vicinity of grain boundary.
at matrix to be substituted by conventional multi-system slip [21,25]. Consequently, grain boundary fracture at and beyond peak aging will be then determined by strain localization in the soft precipitate-free zone. At the same time, overaging of matrix precipitates decreases yield strength of matrix as a result of which matrix deformation incorporates into overall deformation process which is likely to increase tensile ductility and the proportion of intragranular fracture. Further, grain boundary precipitates coarsen during isothermal aging and, hence, the area fraction of grain boundary precipitates decreases by increasing of aging time. It has been well known that the proportion of intragranular fracture and fracture stress increase with decreasing the area fraction of grain boundary precipitates [30]. Furthermore, the preliminary planar array of plate-shaped grain boundary precipitates is modified to spherical morphologies at later stages of aging by virtue of a grain boundary diffusionassisted coarsening reaction, which is expected to decrease stress concentration and, subsequently, increase critical stress or strain for microcrack initiation at precipitate-free zone in the vicinity of grain boundary precipitates. Simultaneously, precipitate-free zone increases in width by increasing aging time. Increasing of precipitate-free zone width leads to decreasing of grain boundary fracture stress if the width of zones exceeds considerably the interparticle spacing of matrix precipitates [31]. Tensile ductil-
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ity also increases slightly as the precipitate-free zone increases in width [30]. In the present steel, width of precipitate-free zone is considerably larger than the spacing of matrix precipitates in all aging times studied. Therefore, presence of precipitate-free zone at grain boundaries of the present steel should be detrimental on the grain boundary fracture of the present steel, tough it may incorporate marginally into tensile ductility at prolonged aging. During further aging beyond peak hardness, overaging of matrix precipitates, coarsening of grain boundary precipitates associated with morphological modification and increasing of precipitate-free zone width occur concurrently. Diffusional transformation into thermodynamically stable fcc austenite also takes place mainly at grain and lath boundaries [17]. A combination of the aforementioned evolutions is expected to increase tensile ductility by increasing of aging time. Nevertheless, tensile ductility still remains zero at intermediate stages up to a few hours. Therefore, it seems that activation of matrix planar slip and coarse grain boundary precipitates in association with precipitate-free zone offset the beneficial microstructural evolutions to improve ductility. After so prolonged aging for larger than 1 day, austenite becomes predominant phase which incorporates to deformation effectively and, consequently, decreases fracture stress and raises ductility by arresting the propagating cracks. Lee et al. [12] attributed the grain boundary fracture of an Fe–10Ni–5Mn maraging steel to closely spaced grain boundary precipitates for which stress–strain incompatibilities at grain boundaries due to intermetallic precipitates could be assumed as an operative mechanism. However, in the present steel grain boundary precipitation was recognized in association with precipitate-free zone. Furthermore, synergistically detrimental effects of inhomogeneous matrix slip and grain boundary weakening were proposed to realize catastrophic grain boundary failure in the early stages of aging. However, more studies are necessary to clarify matrix deformation characteristics and to prohibit grain boundary precipitation reactions of Fe–Ni–Mn steels to suppress their intergranular failure. 5. Conclusions Evolution of the characteristics of grain boundary fracture in correlation with precipitation reactions during isothermal aging at 753 K was studied in an Fe–10Ni–7Mn maraging steel. Main conclusions drawn were as follows: 1. The Fe–10Ni–7Mn maraging steel shows substantial hardening but suffers from premature intergranular failure after aging. 2. Fracture stress turns to increase at initial stages of aging, which unexpectedly decreases after short aging. After a minimum at short aging, fracture stress increases at intermediate stages of aging and eventually decreases at later stages of aging. Tensile elongation decreases drastically at early stages of aging to zero which resumes to a few pct after prolonged aging. 3. A microscopically ductile grain boundary fracture was identified in the aged steel which seems macroscopically brittle
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from the premature onset of fracture and zero tensile elongation. 4. Heterogeneous grain boundary precipitation occurs during isothermal aging in association with solute-depleted precipitate-free zone, realizing a network of weakened material along grain boundaries in the bulk. 5. Inhomogeneous matrix deformation in the form of planar slip bands was found to raise strain localization at precipitate-free zone which reduces critical strain for crack initiation at grain boundary. 6. After prolonged aging, overaging of matrix precipitates, morphological modification of grain boundaries, growth of precipitate-free zone in width increase tensile ductility in association with isothermally transformed austenite particles. References [1] M. Tanaka, T. Suzuki, M. Yodogawa, Bull. Tokyo Inst. Tech. 82 (1967) 79–90. [2] M. Tanaka, J. Yamamoto, Toward Improved Ductility and Toughness, Climax Mo. Dev. Co, Japan, 1971, pp. 195–206. [3] D.R. Squires, E.A. Wilson, Metall. Trans. 3 (1972) 575–581. [4] H.C. Feng, E.A. Wilson, C.J. McMahon Jr., 3rd Int. Conf. on the Strength of Metals and Steels, Cambridge, England, 1973, pp. 129–133. [5] N.H. Heo, H.C. Lee, Metall. Trans. 27A (1996) 1015–1020. [6] N.H. Heo, Acta Mater. 44 (1996) 3015–3023. [7] N.H. Heo, Scripta Mater. 34 (1996) 1517–1522. [8] H. Suto, T. Murakami, Trans. JIM 20 (1979) 365–370.
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