Fracture Toughness of CF8 Castings at Four Kelvin

The first fracture toughness measurements for CF8 stainless steel castings in liquid helium at 4 K are reported. Single-phase (austenite) and duplex (austenite + δ-ferrite) castings were tested. On the basis of estimates from J-integral data, the plane-strain fracture toughness (K_lc) of castings containing 3.2 to 14.5 pct δ-ferrite ranged from 84 to 179 MPa · m^l/2 at 4 K. In contrast, a fully austenitic casting (0 pct δ-ferrite) exhibited a K_lc, value of 331 MPa · m^l/2, which is nearly equivalent to the toughness of a wrought AISI 304 stainless steel of a similar strength. Light and scanning electron microscopy studies indicate that the inferior toughness of castings containing δ-ferrite may be attributed to the brittleness of this body-centered-cubic phase at cryogenic temperatures and its distribution in the microstructure. The relative stability of the austenitic phase with respect to martensitic phase transformation may also play a significant role.     

Effect of Retained Austenite on the Fracture Toughness of Quenching and Partitioning (Q&P)-Treated Sheet Steels

Fracture toughness K _IC was measured by double edge-notched tension (DENT) specimens with fatigue precracks on quenching and partitioning (Q&P)-treated high-strength (ultimate tensile strength [UTS] superior to 1200 MPa) sheet steels consisting of 4 to 10 vol pct of retained austenite. Crack extension force, G _IC, evaluated from the measured K _IC, is used to analyze the role of retained austenite in different fracture behavior. Meanwhile, G _IC is deduced by a constructed model based on energy absorption by martensite transformation (MT) behavior of retained austenite in Q&P-treated steels. The tendency of the change of two results is in good agreement. The Q&P-treated steel, quenched at 573 K (300 °C), then partitioned at 573 K (300 °C), holding for 60 seconds, has a fracture toughness of 74.1 MPa·m^1/2, which is 32 pct higher than quenching and tempering steel (55.9 MPa·m^1/2), and 16 pct higher than quenching and austempering (QAT) steel (63.8 MPa·m^1/2). MT is found to occur preferentially at the tips of extension cracks on less stable retained austenite, which further improves the toughness of Q&P steels; on the contrary, the MT that occurs at more stable retained austenite has a detrimental effect on toughness.     

The effects of heat treatment on fracture toughness and fatigue crack growth Rates in 440C and BG42 steels

The fatigue crack growth rates,da/dN, and the fracture toughness, K_Ic have been measured in two high-carbon martensitic stainless steels, 440C and BG42. Variations in the retained austenite contents were achieved by using combinations of austenitizing temperatures, refrigeration cycles, and tempering temperatures. In nonrefrigerated 440C tempered at 150 °C, about 10 vol pct retained austenite was transformed to martensite at the fracture surfaces duringK _Ic testing, and this strain-induced transformation contributed significantly to the fracture toughness. The strain-induced transformation was progressively less as the tempering temperature was raised to 450 °C, and at the secondary hardening peak, 500 °C, strain-induced transformation was not observed. In nonrefrigerated 440C austenitized at 1065 °C,K _Ic had a peak value of 30 MPa m^1/2 on tempering at 150 °C and a minimum of 18 MPa m^1/2 on tempering at 500 °C. Refrigerated 440C retained about 5 pct austenite, and did not exhibit strain-induced transformation at the fracture surfaces for any tempering temperature. TheK _Ic values for corresponding tempering temperatures up to the secondary peak in refrigerated steels were consistently lower than in nonrefrigerated steels. All of the BG42 specimens were refrigerated and double or quadruple tempered in the secondary hardening region; theK _Ic values were 16 to 18 MPa m^1/2 at the secondary peak. Tempered martensite embrittlement (TME) was observed in both refrigerated and nonrefrigerated 440C, and it was shown that austenite transformation does not play a role in the TME mechanism in this steel. Fatigue crack propagation rates in 440C in the power law regime were the same for refrigerated and nonrefrigerated steels and were relatively insensitive to tempering temperatures up to 500 °C. Above the secondary peak, however, the fatigue crack growth rates exhibited consistently lower values, and this was a consequence of the tempering of the martensite and the lower hardness. Nonrefrigerated steels showed slightly higher threshold values, ΔK_th, and this was ascribed to the development of compressive residual stresses and increased surface roughening in steels which exhibit a strain-induced martensitic transformation.     

Fracture and fractography of metastable austenites

External test variables such as rate and temperature, and changes in alloy composition are shown to have a number of effects on the fracture of high-strength, metastable austenitic steels. One rate-dependent phenomenon is an unusual fracture mode transition wherein a flat mode changes to a shear mode when the amount of transformation product in the vicinity of the crack tip is reduced by adiabatic heating. The point at which this happens in any one test is dependent upon the velocity of the slowly growing crack which in turn is dependent upon the crosshead rate. Because of this rate effect, the plane stress fracture toughness decreases by as much as 30 pct at higher crosshead rates. Fractographically, it was ascertained that at room temperature, both phases failed in a ductile manner, but at ?196°C, martensite containing greater than about 0.27 wt pct C would cleave. This resulted in a “ductile-brittle” transition in metastable austenites at ?196°C as a function of carbon content. Other compositional variations change the austenite stability which controls the amount of strain-induced marteniste occurring at the crack tip. It is shown that a plane stress fracture toughness (K _C) approaching 500,000 psi-in.^1/2 may be achieved by decreasing the stability of the austenite. The variation ofK _c with austenite stability agrees qualitatively with a theoretical model for the invariant shear contribution to the fracture toughness of metastable austenites.     

Effect of overaging on mechanical properties and stress corrosion cracking of HY-180M steel

The tensile properties, fracture toughness and stress corrosion cracking (SCC) behavior of HY-180 M steel at 22 °C were studied after final 5 h overaging treatments >510 ≤650 °C. SCC tests were conducted for 1000 h with compact tension specimens in aqueous 3.5 pct NaCl solutions at a noble (anodic) potential of −0.28 V_SHE ( −0.48 V_Ag/AgC1) and a cathodic protection potential of −0.80 V_SHE (−1.0 V_Ag/AgC1). The SCC resistance improved at aging temperatures >565 °C, the most significant improvement being at −0.80 V_She, especially after 650 ° aging whereK _ISCC was raised to at least 110 MPa · m^1/2. However, this was at the expense of mechanical properties. Provided low crack propagation rates of ∼3 X 10⁻¹¹ m/s at −0.80V _SHEmay be tolerated, the best compromise between strength, toughness, and SCC resistance was obtained after 594 °C aging. Under these conditions, stress intensities as high as ∼ 110 MPa · m^1/2 can be used, with a yield strength of ∼ 1150 MPa and fracture toughness of ∼ 170 MPa · m^1/2. The retained austenite content after aging increased with aging temperature up to 25 pct by vol at 650 °C. It appeared to correlate with improved SCC resistance, but other microstructural effects associated with aging may be involved. Formerly Research Associate with theDepartment of Metallurgical Engineering , University of BritishColumbia     

Microstructure and fracture of 52100 steel

The fracture behavior of 52100 steel hardened and tempered to R_C62 has been investigated as a function of austenitizing over the temperature range from 800 to 1100°C. Specimens were homogenized at 1150°C and either furnace cooled or isothermally transformed at 580°C to produce a pearlitic microstructure prior to austenitizing for hardening. Furnace-cooled specimens developed a proeutectoid carbide network that did not dissolve during subsequent austenitizing below A_cm . The residual proeutectoid carbides and the carbide-free martensite-austenite structure between them controlled fracture and produced K_IC of 19 MPa \ m^1/2, the highest determined in this investigation. The specimens isothermally transformed prior to austenitizing below A_cm produced a microstructure of fine spherical carbides dispersed throughout a fine martensitic matrix and did not contain residual proeutectoid carbides. The transgranular fracture of the latter specimens by microvoid coalescence around the closely spaced spherical carbides resulted in the lowest values of fracture toughness, 14 to 16 MPa\ m^1/2, determined in these experiments. Austenitizing above A_cm caused solution of all carbides, a gradual coarsening of the austenitic grain size, a transition to plate martensite, and an increase in retained austenite. Fracture toughness increased slightly with increasing austenitizing temperature above A_cm despite the fact that fracture propagated primarily along the austenitic grain boundaries. The improved fracture toughness, verified by scanning electron microscopy of the fatigue crack-overload fracture interface, is believed to be caused in part by transgranular crack propagation during the first stages of crack extension that are most important in determining K_1C.     

The effect of retained austenite on the fracture toughness of high-speed steels

The aim of this work was to find the quantitative dependences between fracture toughness K_lc and the volume fraction of retained austenite in the matrix of quenched high-speed steels. The tests were carried out on three model alloys of a different content quotient of Mo: W which, after quenching, were gradually supercooled up to ? 196°C and then tempered at 450°C. Also the measurements of the content of retained austenite in the vicinity of the surface of a sample fracture were carried out. It was determined that after tempering at 450°C the fracture toughness of the matrix of high-speed steels is directly proportional to the content of retained austenite in it. Every 1 % by volume of retained austenite increases the fracture toughness K_lc of the matrix by about 5%, despite the fact that most probably it is completely transformed into fresh martensite in front of a propagating crack. Higher fracture toughness of the matrix of high-speed steels rich in molybdenum should be explained exlusively by a larger content of retained austenite. Transformations in the martensitic part of the matrix of the alloys richer in molybdenum clearly reduce the advantageous effect of retained austenite on this steel feature.     

A dislocation shielding prediction of the toughness transition during cleavage of semibrittle crystals

An interconnected set of observations assesses current equilibrium models of the ductile-brittle-transition temperature (DBTT). This involvesin situ transmission electron microscopy (TEM) studies of crack-tip dislocations in single and polycrystals and bulk fracture toughness tests at various temperatures. Beyond K_I values of 8 MPa · m^1/2 in both iron-base single and polycrystals, large numbers of redundant dislocations are created, as postulated recently by Weertman. [38] Still, the necessary shielding dislocations, as required by equilibrium, can be detected at values as high as 20 and 40 MPa · m^1/2 byex situ TEM and electron channeling, respectively. In addition, the close approach of dislocations to the crack tip in some of the studies, as opposed to others, suggests that large dislocation free zones (DFZ) are a thin-film artifact. However, a failure criterion based partly on the Rice-Thomson model’21 is both consistent with the absence of a large DFZ and observed fracture toughness variations with test temperature. It is emphasized that this toughness transition is entirely in the semibrittle regime where cleavage is the failure mode. Nevertheless,K _lc values increase from 3 to 60 MPa·m^1/2 with an increase in test temperature. This article is based on a presentation made in the symposium “Quasi-Brittle Fracture” presented during the TMS fall meeting, Cincinnati, OH, October 21–24, 1991, under the auspices of the TMS Mechanical Metallurgy Committee and the ASM/MSD Flow and Fracture Committee.     

An investigation of the high-temperature and solidification microstructures of PH 13-8 Mo stainless steel

Differential thermal analysis (DTA), high-temperature water-quench (WQ) experiments, and optical and electron microscopy were used to establish the near-solidus and solidification microstructures in PH 13-8 Mo. On heating at a rate of 0. 33 °C/s, this alloy begins to transform from austenite to δ-ferrite at ≈1350 °C. Transformation is complete by ≈1435 °C. The solidus is reached at ≈1447 °C, and the liquidus is ≈1493 °C. On cooling from the liquid state at a rate of 0. 33 °C/s, solidification is completed as δ-ferrite with subsequent transformation to austenite beginning in the solid state at ≈1364 °C. Insufficient time at temperature is available for complete transformation and the resulting room-temperature microstructure consists of matrix martensite (derived from the shear decomposition of the austenite) and residual δ-ferrite. The residual δ-ferrite in the DTA sample is enriched in Cr (≈16 wt pct), Mo (≈4 wt pct), and Al (≈1. 5 wt pct) and depleted in Ni (≈4 wt pct) relative to the martensite (≈12. 5 wt pct Cr, ≈2 wt pct Mo, ≈1 wt pct Al, ≈9 wt pct Ni). Solid-state transformation of δσ γ was found to be quench-rate sensitive with large grain, fully ferritic microstructures undergoing a massive transformation as a result of water quenching, while a diffusionally controlled Widmanstätten structure was produced in air-cooled samples.     

Orientation relationships among M23C6, M6C,and austenite in an Fe-Mn-Al-Mo-C alloy

Both M₂₃C₆ and M_i6C carbides were observed to precipitate within the austenite phase in an Fe-24.6 pct Mn-6.6 pct Al-3.1 pct Mo-1.0 pct alloy after being quenched from 1200 °C and aged at 700 °C. By means of transmission electron microscopy and diffraction techniques, the orientation relationships among M₂₃C₆, M₆C, and the austenite phase were determined as follows: {fx567-1} The present result of the orientation relationship between M₆C and the austenite phase is in disagreement with that reported by Maziasz^[14] for M₆C in an austenitic stainless steel.     

Dependence of fracture toughness of austempered ductile iron on austempering temperature

Ductile cast iron samples were austenitized at 927 °C and subsequently austempered for 30 minutes, 1 hour, and 2 hours at 260 °C, 288 °C, 316 °C, 343 °C, 371 °C, and 399 °C. These were subjected to a plane strain fracture toughness test. Fracture toughness was found to initially increase with austempering temperature, reach a maximum, and then decrease with further rise in temperature. The results of the fracture toughness study and fractographic examination were correlated with microstructural features such as bainite morphology, the volume fraction of retained austenite, and its carbon content. It was found that fracture toughness was maximized when the microstructure consisted of lower bainite with about 30 vol pct retained austenite containing more than 1.8 wt pct carbon. A theoretical model was developed, which could explain the observed variation in fracture toughness with austempering temperature in terms of microstructural features such as the width of the ferrite blades and retained austenite content. A plot of K _IC ² against σ _y (X _γ, C _γ)^1/2 resulted in a straight line, as predicted by the model.     

The fracture characteristics of Al-9Ti/SiC p metal matrix composites

A fracture mechanics approach was used to determine the plane strain fracture toughness (K _IC) of a mechanically alloyed Al-9Ti 20 vol pct cobalt sol-gel-coated SiC particle-reinforced composite. Processing defects consisting of clumped SiC particulate, bonded by the sol-gel, initiated failure in tensile tests. The defects were measured and the fracture toughness was calculated using the Irwin relation. The value ofK _IC for the as-received material was determined to be equal to 4.7 MPa·m^1/2 at room temperature. Annealing the material for 120 hours and 400 hours at 500 °C increased the fracture toughness. This can be attributed to coarsening of an Al₃Ti strengthening phase. Tensile tests conducted at 200 °C show thatK _IC decreases at that temperature for each annealing condition. The sensitivity to the presence of the defects is greatest for samples annealed at 500 °C for 120 hours. The effect of the defects on the failure mechanism of the composite material as a function of temperature was determined. At room temperature, the Co/SiC processing defects provide low-energy paths for crack propagation; at 500 °C, the defects serve as void nucleation sites.     

The effects of titanium additions on AF1410 ultra-high-strength steel

The mechanical properties and microstructure of two heats of AF1410 steel were compared. The first heat, heat 811, contained a titanium addition of 0.02 wt pct, while the second heat, heat 91, contained no titanium, manganese, or other strong sulfide formers. The sulfur in heat 811 was gettered as titanium carbosulfide, while in heat 91 the sulfides were chromium sulfide. The toughness of heat 811 was found to be much enhanced compared to heat 91, with Charpy impact energies of 176 J and 79 J and K_IC fracture toughness values of 235 MPa.m^1/2 and 170 MPa.m^1/2, respectively. This significant difference in fracture toughness is attributed to the fact that titanium carbosulfide particles are more resistant to void nucleation than the chromium sulfide particles, which appear to nucleate voids at the onset of plastic strain. In addition to altering the sulfide particle type, the titanium addition also results in the presence of undissolved MC carbides in the titanium-modified steel in addition to the M₂C carbides found in heat 91. These carbides act as grain growth inhibitors, resulting in a finer prior austenite grain size and martensite packet size in heat 811.     

Nanoscale Analyses of High-Nickel Concentration Martensitic High-Strength Steels

Austenite reversion in martensitic steels is known to improve fracture toughness. This research focuses on characterizing mechanical properties and the microstructure of low-carbon, high-nickel steels containing 4.5 and 10 wt pct Ni after a QLT-type austenite reversion heat treatment: first, martensite is formed by quenching (Q) from a temperature in the single-phase austenite field, then austenite is precipitated by annealing in the upper part of the intercritical region in a lamellarization step (L), followed by a tempering (T) step at lower temperatures. For the 10 wt pct Ni steel, the tensile strength after the QLT heat treatment is 910 MPa (132 ksi) at 293 K (20 °C), and the Charpy V-notch impact toughness is 144 J (106 ft-lb) at 188.8 K (?84.4 °C, ?120 °F). For the 4.5 wt pct Ni steel, the tensile strength is 731 MPa (106 ksi) at 293 K (20 °C) and the impact toughness is 209 J (154 ft-lb) at 188.8 K (?84.4 °C, ?120 °F). Light optical microscopy, scanning electron and transmission electron microscopies, synchrotron X-ray diffraction, and local-electrode atom-probe tomography (APT) are utilized to determine the morphologies, volume fractions, and local chemical compositions of the precipitated phases with sub-nanometer spatial resolution. The austenite lamellae are up to 200 nm in thickness, and up to several micrometers in length. In addition to the expected partitioning of Ni to austenite, APT reveals a substantial segregation of Ni at the austenite/martensite interface with concentration maxima of 10 and 23 wt pct Ni for the austenite lamellae in the 4.5 and 10 wt pct Ni steels, respectively. Copper-rich and M₂C-type metal carbide precipitates were detected both at the austenite/martensite interface and within the bulk of the austenite lamellae. Thermodynamic phase stability, equilibrium compositions, and volume fractions are discussed in the context of Thermo-Calc calculations.     

Stress corrosion cracking of stainless steels in NaCl solutions

The metallurgical influences on the stress corrosion resistance of many commercial stainless steels have been studied using the fracture mechanics approach. The straight-chromium ferritic stainless steels, two-phase ferritic-austenitic stainless steels and high-nickel solid solutions (like alloys 800 and 600) investigated are all fully resistant to stress corrosion cracking at stress intensity (K₁) levels ≤ MN • m-^3/2 in 22 pct NaCl solutions at 105 °C. Martensitic stainless steels, austenitic stainless steels and precipitation hardened superalloys, all with about 18 pct chromium, may be highly susceptible to stress corrosion cracking, depending on heat treatment and other alloying elements. Molybdenum additions improve the stress corrosion cracking resistance of austenitic stainless steels significantly. The fracture mechanics approach to stress corrosion testing of stainless steels yields results which are consistent with both the service experience and the results from testing with smooth specimens. In particular, the well known “Copson curve” is reproduced by plotting the stress corrosion threshold stress intensity (AT_ISCC) vs the nickel content of stainless steels with about 18 pct chromium. Formerly with the BBC Brown Boveri Company, Baden, Switzerland     

Effect of cooling after welding on microstructure and mechanical properties of 12 Pct Cr steel weld metals

The microstructure of three 12 pct cr steel weld metals with different nickel and nitrogen contents was studied in as-welded condition and after postweld heat treatment with and without intercooling. Tensile strength and impact toughness of the weld metals were investigated in different postweld heat treatment conditions. In weld metals heat treated without intercooling, austenite decomposed by a eutectoid reaction that resulted in M₂₃C₆ aggregates around retained δ-ferrite. Two morphologies of M₂N and MN precipitates were found in a low-dislocation α-ferrite. It was concluded that these phases were also transformed from austenite. In weld metals heat treated with intercooling, M₂₃C₆ precipitates were smaller and more homogeneously distributed. Different MN precipitates were found in the tempered martensite. The fracture mode of the weld metals at room temperature was mainly transgranular cleavage with some fibrous fracture. Intercooling treatment improved Charpy impact toughness of the 12 pct Cr steel weld metals substantially. It was found that the important microstructural factors affecting the impact toughness of the weld metals which were heat treated without intercooling were the sizes of the α-ferrite grains, nonmetallic inclusions, and M₂₃C₆ aggregates. For the weld metals heat treated with intercooling, the factors which affect the toughness of the weld metals were the sizes of martensite packets and nonmetallic inclusions.     

Effects of Strength and Microstructure on Hydrogen-Assisted Crack Propagation in 22Cr-13Ni-5Mn Stainless Steel Forgings

The objective of this study was to evaluate the effects of hydrogen on the fracture toughness and fracture mechanisms for the nitrogen-strengthened, austenitic stainless steel 22Cr-13Ni-5Mn, an alloy with potential value in high-pressure hydrogen containment components. The fracture initiation toughness and crack-growth resistance were measured before and after thermal precharging with hydrogen and as a function of crack-growth orientation and material strength. The effects of crack-growth orientation and material strength dominated over the effect of hydrogen exposure. The former two variables caused changes in fracture initiation toughness of up to 400 pct, while dissolved hydrogen resulted in only modest decreases in fracture initiation toughness of 20 to 40 pct. Coarse Z-phase (CrNbN) particles aligned in bands governed the measured fracture toughness and observed fracture mode. Fracture progressed by void nucleation and growth in the Z-phase bands, forming microcracks that ultimately linked through the remaining austenite matrix. Crack-growth orientation, material strength, and hydrogen exposure affected the nucleation and growth of voids in the Z-phase bands and the subsequent linking of microcracks. Control or elimination of the coarse, banded Z phase would likely enhance the fracture resistance of this alloy.     

Fracture and fatigue in M-50 and 18-4-1 high speed steels

The plane strain fracture toughness values,K _Ic , the fatigue crack growth rates and the tensile properties of M-50 and 18-4-1 high speed steels have been measured as a function of tempering temperature. The M-50 was a vacuum arc remelted grade (VIM-VAR) and the 18-4-1 an electroslag grade, and both are used in mainshaft gas turbine bearings. At the usual hardness for bearings, Rockwell C 62, 18-4-1 exhibited a slightly higher fracture toughness (21 MPa·m^1/2) than M-50 (18 MPa·m^1/2). The fatigue crack growth rates were very similar, and in the slow growth region followed the usual power law,dC/dN=(ΔK) ^m withm=3 to 4. The crack propagation rates were still significant at values as low as ΔK=5 MPa·m^1/2. SEM studies of the fracture surfaces showed complex transgranular fracture paths for both steels. The tensile strengths and the elongations of M-50 were somewhat higher than the corresponding values for 18-4-1 but the yield strengths of the two steels were similar. The microstructures of these steels were markedly different, with M-50 exhibiting 2.6 vol pct undissolved carbides and the 18-4-1 showing 15.2 vol pct carbides, but the fatigue and fracture behaviors were similar.     

Microstructural change during superplastic deformation of δ-ferrite/austenite duplex stainless steel

Superplasticity of a 25 pct Cr-6.5 pct Ni-3 pct Mo-0.14 pct N δ/γ duplex stainless steel has been studied with particular emphasis on the microstructural change during deformation. Two large superplastic elongations are obtained at temperatures around 1323 K in δ/γ duplex phase region and 1173 K where σ phase particles precipitate dynamically at a strain rate of ~10^?3 s^?1. During deformation in the higher temperature region, fine Widmanstätten γ particles coarsen and coarse γ grains formed during the prior treatments are broken into spherical particles, resulting in a homogeneous dispersion of γ particles within the σ-ferrite matrix. The dynamic recrystallization of soft σ-ferrite matrix occurs locally in the region where the strain reaches some critical value, and the final microstructure consists of equiaxed σ and γ grains. In the case of lower temperature deformation, a eutectoid decomposition of δ-ferrite into γ and σ phases occurs. The relatively soft γ grains which are severely deformed by hard σ particles recrystallize dynamically, and these processes lead to the γ/σ equiaxed duplex structure. The extremely large superplasticity of this alloy can mainly be explained in terms of the above microstructural change during deformation.