It is clear that the transition-temperature shift when a steel is temper-embrittled depends strongly on the steel's hardness, which in turn is a function of the microstructure. Four main types of low-temperature intergranular fracture will be briefly reviewed, demonstrating mainly the effects of minor impurity elements.
These are characterized by a fracture path which follows prior austenite grain-boundaries, but which, in detail, is composed of very fine micro-voids, centres around small sulfide inclusions located at the boundaries (Figure 1). Macroscopically, the fracture surface exhibits large, dull facets. Such fractures can be obtained in high quality, low-sulphur, forging steels, which are austenitized at temperatures sufficiently high (greater than approx. 1200°C) for virtually all the sulfur to be taken into solid solution. Then, for some rates of cooling, iron-, manganese- or chromium-sulfides precipitate preferentially as a fine distribution at the austenite boundaries producing closely-spaced void-nucleating centres, which are responsible for the intergranular ductile fracture path.
Figure 1: Overheated fracture surface in low-alloy steel. (x 150) (x 5/7)
The sulfide network persists on tempering, even at temperatures up to 700°C. In high-strength, lightly tempered, forging steels, the presence of intergranular facets is not associated with particularly low values of toughness, because the steels have to be low in sulfur for the effect to be observed. Indeed, in sharply pre-cracked test-pieces, the conventional fracture toughness of overheated steel can be high, because it is far more likely for the tip of the fatigue pre-crack to be located in the interiors of sulfide-free grains than along sulfide-decorated boundaries. In blunt notched specimens, on the other hand, the apparent toughness is substantially reduced, because an easy fracture path (a decorated boundary) is always present somewhere in the high strain region at the root of the notch.
The many experiments have shown that the solubility of phosphorus in austenite varies markedly with temperature and that its segregation to grain boundaries at relatively low temperatures (<1050°C) is responsible for the production of brittle intergranular fractures in as-quenched steels. At high temperatures (>1200°C) very little phosphorus is segregated at boundaries and the low-temperature fracture mode of steels quenched from such temperatures is transgranular cleavage. The fracture surface of forging steels quenched from 1300°C and broken at 77 K is shown in Figure. 2: due to the very coarse grain size and martenzitic packet size, the deviation of the cleavage crack through individual martenzitic laths can be clearly recognized. The variation of the amount of phosphorus segregated at grain boundaries, as a function of austenitizing temperature, is one possible reason for the observed increase in KIC values in as-quenched steels for austenitizing temperatures greater than 1150°C.
Figure 2: Transgranular cleavage in 1300°C. (X 190) (x5/7)
It is commonly observed that the room-temperature notched-impact energy or fracture toughness of quenched-and-tempered steels shows a minimum for tempering temperatures in the range 300°C - 400°C. This is known as "350°C" or "500°F" Embrittlement. Sometimes, in KIC tests, the minimum is not clear for room temperature tests, but becomes apparent on testing at sub-zero temperatures. The modes of fracture associated with 350°C Embrittlement may be transgranular cleavage, intergranular cleavage or even fibrous (micro-void-coalescence). In work carried out, using a simple Fe-0.6%C alloy, in which the mode was transgranular cleavage throughout, it was argued that the 350°C minimum occurred, because the microstructure at 350°C (the bi-modal distribution of cementite, Figures 2 and 3 /see Part One/ showed the worst combination of high matrix yield strength (due to the lath walls and fine, intra-lath carbides) and coarse crack nuclei (inter-lath, inter-packet or grain-boundary carbides). At lower tempering temperatures, the crack nuclei were less coarse: at higher tempering temperatures, the matrix was softer, and so the tensile stress available to propagate cleavage cracks was lower. A similar mechanism can be invoked to explain intergranular 350°C Embrittlement fractures, with the further complication that concentration of impurity elements at prior austenite grain boundaries (either by segregation or by "carbide rejection" occurs during tempering. This may be demonstrated by examining the fracture surfaces of specimens broken at 77 K: a change from transgranular cleavage to intergranular cleavage and back to transgranular cleavage is observed, as the tempering temperature is increased. Equilibrium segregation of phosphorus during tempering has been used to explain a form of stress-relief-cracking in 2Cr-1Mo steels. The cause of a toughness minimum at 350°C for fully fibrous fractures is not clear, but could again be associated with the long, plate-like carbides which are present at 350°C. The strain required to crack, and hence produce a void nucleus in a plate carbide is much less than that for spheroidal carbides.
In low strength steel, above the transition temperature, micro-cracks do not propagate as sharp cracks, but become blunted. With further tensile strain, blunted cracks in grain-boundary carbides can open up as voids. Fracture in a tensile test proceeds by a combination of external necking (which would lead to separation at a point, or chisel edge in a pure iron specimen) and internal necking, leading to the coalescence of voids formed around second-phase particles in the interior of the specimen. Voids initiate most easily around non-metallic inclusions, particularly around MnS which, unlike oxide or silicate inclusions, contracts on cooling more than the surrounding iron matrix. Small plastic strains are required to initiate cracks/voids in plate-like carbides, either by the action of dislocation pile-ups or by fibre-loading, if the plates are roughly parallel to the tensile axis. Larger plastic strains are needed to initiate voids around spheroidal carbides. In a later section, the argument will be developed that this is a consequence of the number of dislocations tangled around the particles and stressing the carbide/matrix interface.
In marageing steels, intermetallics are used as hardening particles instead of carbides and it is arguable that the superior toughness of marageing steels is attributable to the fact that the intermetallics do not crack easily and that the intermetallic/matrix interface is strong. In fact, deterioration of toughness in marageing steels is produced by unfavourable distributions of small amounts of titanium carbo-nitrides. More generally, the toughness of steels above the transition temperature is controlled primarily by the volume fraction and distribution of non-metallic inclusions. Matrix properties become increasingly significant at higher strength levels in high-quality, low sulphur forging steels, or if a low strength steel has been heavily cold-worked. Low fracture displacements are associated with closely-spaced inclusions and lower-than-normal toughnesses may be associated with short-transverse (ST) orientations in wrought plates, with inter-dendritic MnS inclusions in castings, or with finely-dispersed deoxidation products in weldments.