Austenite-martensite Interface Research Articles

The pressure dependence of the austenite start temperature in iron-nickel base alloys

The effect of hydrostatic pressure up to 20 kbar on the austenite start temperature, A s of an Fe-30.3 wt.% Ni alloy has been measured in detail. An exceptionally large pressure coefficient of ≥ 30°C/kbar is observed up to 2.3 kbar; from 2.3 to 6 kbar A s is observed to increase with pressure, while from C to 20 kbar A s is observed to decrease at a rate of −1.9°C/kbar. A model based on the difference in compressibilities of the martensite and austenite phases is proposed to explain these results. Cold working by rolling prior to the application of hydrostatic pressure shifts the general features of the non-cold worked A s vs pressure curve to higher pressure and to higher temperature. It is argued that these observations give very strong evidence that martensite reversal initiates at the martensite-austenite interface. Similar studies were also made on Fe-29.1 wt.% Ni-1 wt.% Co and Fe-30.4 wt.% Xi-2 wt.°% Mn ternary alloys. These results indicate that the character of the martensite reversal (athermal vs. isothermal) may affect the dependence of A s on hydrostatic pressure.

The effect of plastic deformation on the martensite-to-austenite transition in an iron-nickel alloy

The effect of plastic deformation introduced by rolling at room temperature on the austenite start temperature of an Fe-30.3 wt pct Ni-0.005 wt pct C alloy has been determined. The austenite start temperature increases monotonically with deformation. Microhardness measurements show that the austenite start temperature increases with the yield strength of the martensite. The temperature at which martensite reversal initiates is not affected by the amount of martensite present, and, therefore, is not dependent on the martensite plate size. It is suggested that the reverse martensite transformation initiates at the martensite-austenite interface and is controlled by interface propagation.

The morphology and substructure of highly tetragonal martensites in Fe-7 pct Al−C steels

This investigation reports on an optical and electron microscope study of bct martensite formed in Fe-7 pct Al-1.5 pct C and Fe-7 pct Al-2.0 pct C alloys. In each case the martensite is plate-like containing\((112)[\bar 1\bar 11]\) transformation twins 100 to 200A in width. The particular twin plane variant\((112)[\bar 1\bar 11]\) corresponds to the martensite habit plane variant (3, 15, 10)F, which is predicted by the crystallography theory. The twins are uniformly spaced and extend completely from one martensite-austenite interface to the other as would be theoretically expected. The martensite plates are ideally lenticular in the 2 pct C alloy but those in the 1.5 pct C alloy frequently exhibit irregular interfaces which are attributed to impingement effects. All observations are in accordance with the phenomenological crystallography theory as applied to ferrous martensites with a {3, 15, 10}F habit plane.

A new observation on the mechanism of martensitig transformations

In an Fe-Ni-P alloy the influence of incoherent particles in the austenite on subsequent transformation to martensite is investigated. It is observed that small cone-shaped regions of retained austenite are left like “shadows” behind the particles after the transformation interface has passed. The geometrical and crystallographic features of these “austenite shadows” are determined and analysed by a comparison with the elements of the phenomenological theory of the transformation; these elements have been obtained by numerical calculations from the knowledge of the lattice constants, the orientation relationship and the habit plane. The main result of this analysis is that the direction of the “austenite shadows” is antiparallel to the direction of the macroscopic shear. From this new geometrical conditions are deduced which have to be satisfied by the mechanism responsible for the propagation of the austenite-martensite interface during the transformation. Furthermore it has been observed that the directions of the “austenite shadows” have opposite signs on either side of the plane defined by the midrib. From this it follows that the midrib plane is the plane of initiation of the transformation.

Transmission electron microscopy studies of {225} f martensite in an Fe-8%Cr-1%C alloy

Fully grown {225}, martensite plates in an Fe-8 %Cr-1 %C alloy were examined by means of transmission electron microscopy and diffraction. (112) b twins, (011) b planar defects and (1̄12) b slip traces were observed as internal planar inhomogeneities accompanying the transformation.§ General features of the (112) b twins were considerably different from those found in ferrous martensites which exhibit the {3, 10, 15} f habit plane. The (112) b twins were usually observed only at one austenite-martensite interface and a mid-rib region of high twin density was not found in most plates. In a few instances a mid-rib was observed at the center of plates, in which cases the (112) b transformation twins were bent at this region as well as at the interface regions to which the twins extended. (1̄12) b slip resulted in fragmented (112) b twins in several cases. Internal (011) b planar defects, which are possibly twins, formed on planes parallel with the (111) f planes of the Kurdjumov-Sachs orientation relationship. In addition to the above planar inhomogeneities, lengthy dislocations nearly parallel to (112) b and (011) b were observed in untwinned regions of plates. These internal dislocations were mixed and not of screw type as observed in Fe-Ni martensites. Dislocations observed at the austenite-martensite interface were nearly parallel to [11̄2] f , which lies approximately in the habit plane. These various internal defects are considered in relation to existing crystallographic information and theory.

Austenite–Martensite Interface Dislocations in Stainless Steel

WHEN 18 chromium : 8 nickel austenitic stainless steel is plastically deformed at room and sub-zero temperatures, e (hexagonal closest packing) and α′ (body centred cubic) martensites are formed. The e phase forms as the result of the propagation of Shockley partial dislocations on every second {111} γ plane and appears in the form of thin disks having a {111} γ habit1. The mechanism of the γ → α′ transformation in this steel is not yet entirely understood. Venables1 suggested that α′ martensite nucleates from the e phase and grows by the motion of a parallel grid of screw dislocations such as that postulated by Frank2. Bogers and Burgers3 showed, by geometrical arguments, that the shears necessary for the transformation correspond to ‘well-known’ (that is, Shockley partial) dislocations and also suggested the possibility that the transformation proceeded by way of the motion of a parallel array of screw dislocations. Lagneborg4 found that deformation induced α′ nucleated preferentially at the intersections of two e disks or at the intersections of active slip planes with e disks, and proposed that the martensite-like region existing near the centre of a Shock-ley partial dislocation wras a favourable nucleation site. He did not discuss a mechanism for the subsequent growth of α′. To our knowledge, no direct evidence for the postulated dislocation interfaces has been observed. During a recent investigation of the effects of shock loading on 18 chromium : 8 nickel stainless steel, an experimental observation was made that showed a matrix-martensite interface composed of an array of dislocations. This communication presents this observation and discusses its relevance to the previously postulated transformation mechanisms.

The crystallography and growth of partially-twinned martensite plates in Fe-Ni alloys

Observations of the internal structure, morphology, and surface relief associated with partiallytwinned plates in Fe-Ni alloys containing near 30 wt. % Ni indicate that the growth of such plates involves two stages which are indistinguishable on the basis of the shape change of the transformed regions; however these stages differ in the transformation inhomogeneity and in the growth kinetics. Stage one corresponds to an internally twinned region of a plate and stage two corresponds to an untwinned region. It has been verified that the internal twinning resulting from first-stage growth is the inhomogeneous shear of the phenomenological crystallographic theory; i.e. the observed twinning shear gives rise to an invariant (habit) plane consistent with the observed midrib plane trace. The inhomogeneous shear associated with second-stage growth appears to be slip whose elements are the same as the twinning elements of stage one. The transformation inhomogeneity probably dictates the final martensite morphology since if only stage one is involved in the growth of a (completely-twinned) plate the austenite-martensite interface is planar in the classical sense; however partially-twinned plates do not exhibit planar interfaces and if no twinning occurs at all, plates are not formed. It is suggested that the local temperature rise produced by the (exothermic) transformation effects the transition from stage one to stage two.

Coherent growth of martensite during tempering

It is shown that martensite plates in a 1.5% C 5.0% Ni steel can, during tempering, thicken or grow into bainite. It is suggested that this supports the author's view on bainite formation that the reduction in volumetric strain by the removal of carbon in the form of carbide during the formation of bainite provides part of the additional driving force needed for the coherent growth to proceed. The results also suggest that the austenite-martensite interface remains coherent after the development of a martensite plate has ceased.