ASTM Standard Test Method E 1921-97, “Test Method for the Determination of Reference Temperature, T0, for Ferritic Steels in the Transition Range, addresses determination of T0, a fracture toughness reference temperature for ferritic steels having yield strength ranging from 275 to 825 MPa. E 1921 defines a ferritic steel as: “Carbon and low-alloy steels, and higher alloy steels, with the exception of austenitic stainless steels, martensitic, and precipitation hardened steels. All ferritic steels have body centered cubic crystal structures that display ductile to cleavage transition temperature. This definition is not intended to imply that all of the many possible types of ferritic steels have been verified as being amenable to analysis by this test method.” The equivocation provided by the final sentence was introduced due to lack of direct empirical evidence (i.e., fracture toughness data) demonstrating Master Curve applicability for all ferritic alloys in all heat treatment/irradiation conditions of interest. This question regarding the steels to which E 1921 applies inhibits its widespread application for it suggests that the user should perform some experimental confirmation of Master Curve applicability before it is applied to a new, or previously untested, ferritic steel. Such confirmations are, in many cases, either impractical to perform (due to considerations of time and/or economy) or impossible to perform (due to material unavailability). In this paper we propose an alternative to experimental demonstration to establish the steels to which the Master Curve and, consequently, ASTM Standard Test Method E 1921 applies. Based on dislocation mechanics considerations we demonstrate that the temperature dependency of fracture toughness in the fracture mode transition region depends only on the short-range barriers to dislocation motion established by the lattice structure (body-centered cubic (BCC) in the case of ferritic steels). Other factors that vary with steel composition, heat treatment, and irradiation include grain size/boundaries, point defects, inclusions, precipitates, and dislocation substructures. These all provide long-range barriers to dislocation motion, and so influence the position of the transition curve on the temperature axis (i.e., T0 as determined by E 1921-97), but not its shape. This understanding suggests that the myriad of metallurgical factors that can influence absolute strength and toughness values exert no control over the form of the variation of toughness with temperature in fracture mode transition. Moreover, this understanding provides a theoretical basis to establish, a priori, those steels to which the Master Curve should apply, and those to which it should not. On this basis, the Master Curve should model the transition fracture toughness behavior of all steels having an iron BCC lattice structure (e.g., pearlitic steels, ferritic steels, bainitic steels, and tempered martensitic steels). Conversely, the Master Curve should not apply to untempered martensitic steels, which have a body-centered tetragonal (BCT) lattice structure, or to austenite, which has a FCC structure. We confirm these expectations using experimental strength and toughness data drawn from the literature.
The cleavage fracture model of Wallin et al. [1] suggests that the variation of median and bounding values of fracture toughness with temperature is predictable on a micro-mechanical basis. ASTM has recently adopted a standard [ASTM Standard Test Method for the Determination of Reference Temperature, To, for Ferritic Steels in the Transition Range (E 1921–98)] for characterizing the variation of the fracture toughness of ferritic steels in the transition temperature range that draws heavily on the work of Wallin et al. and on subsequent developments [1–5]. The new standard expresses the median toughness transition for a 1T specimen as: KJc|median=30+70⋅exp[0.019(T−To)] where temperature is measured in degrees Celsius and KJc is measured in MPa√m. Similar equations express the variation of bounding toughness values (e.g., 95% lower bound) with temperature. The numeric coefficients in all of these equations do not depend on the type of steel tested. This invariance suggests that chemistry, heat treatment, and other metallurgical variables are not thought to influence the exponential increase with temperature of the plastic work necessary for crack propagation, the dominance of carbides as the particles that initiate cleavage fracture, or the distribution of carbide particles. The ASTM E 1921 Master Curve represents existing fracture toughness data well for nuclear pressure vessel steels and their weldments [6]. This empirical evidence suggests that the material invariance attributed to the Master Curve coefficients in ASTM E 1921 is at least approximately correct for this class of steel. But the microstructural bounds of applicability for the Master Curve are not clear. In this paper we examine the physical basis for the Wallin et al. Master Curve with the aim of distinguishing the classes of steels to which the methodology applies from those to which it does not. We use this physical understanding to calculate the temperature dependence of the plastic work for ferritic steels to demonstrate theoretical validity of a single “master curve.”
An investigation of the combined effects of slow strain rate testing and cathodic protection of Ni-Cu Alloy K-500 forgings was performed. The J-integral fracture toughness tests were performed in a modified screw-driven Instron testing machine with measured cross head travel speeds of 0.1, 0.013, and 0.001 mm/min (0.004, 0.0005, and 0.000 04 in./min). The IT compact specimens used were modified for use with direct current potential drop (DCPD) measurement of the crack lengths. Specimens tested at 0.1, 0.013, and 0.001 mm/min were in an air environment. Additional specimens were tested at 0.001 mm/min in a 3.5% NaCl solution with an applied cathodic potential of -1.0 V (SCE). The tests were monitored with a desk top computer data acquisition system that recorded load, load-line opening displacement (COD), DCPD, applied cathodic potential, and impressed current. These data were used to construct J-integral resistance (J-R) curves for all the samples. The J-R curves for all the specimens tested in air had appeared similar. The specimens tested in salt water with cathodic protection had significant reduction in J-integral toughness (JIc) to less than half the air value. Examination of the specimens' fracture surfaces revealed that the cathodic protection caused the fracture mode to change from dimpled rupture to intergranular fracture.
An abundance of empirical data supports the use of the Master Curve, as proposed by Wallin, Saario and Törrönen, to describe the fracture toughness transition behavior of ferritic steels, particularly the notion of a curve shape that is invariant with steel microstructure (other than lattice structure). However, nuclear surveillance programs do not always contain samples of the steel that most limits reactor operations, making direct measurement of fracture toughness impossible. This suggests that a purely empirical argument cannot define the limits of applicability of the Master Curve or validate its use for all conditions of interest. In previous papers a microstructural basis for the existence of a single “Master” fracture toughness transition curve for all ferritic steels was established and limits of applicability have begun to be explored from a theoretical viewpoint. These previous papers established that all steels with the same lattice structure and cleavage fracture mechanism should be expected to adhere to transition behavior that can be defined by a single curve shape with variations in microstructure accounting only for a shift in the transition temperature. In this paper we explore the basis for “Master Curve” validity for irradiated steels by exploring how irradiation affects the microstructure and fracture mode and using the Zerilli-Armstrong constitutive model as the basis for predictions of irradiated steel behavior.
It is shown that pre-crack and early stage fatigue damage can be characterized by a new sensor technology, the Meandering Winding Magnetometer (MWM™). This new technology consists of a conformable sensor, the MWM, and associated measurement grids that are model based. Measurements on type 304 stainless steel indicated that damage is readily detectable at 20% of life (N/NF = 0.2) and causes a 1.5% conductivity loss. Near failure the conductivity loss in the crack-free region was approximately 4%. In 2024 aluminum the onset of detectable fatigue damage was observed at approximately 50% of total life. For the probe geometry employed, the conductivity loss in the microcrack region just prior to failure was 7%; in the macrocrack region it reached 13%.
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