NON-EQUILIBRIUM STRUCTURES AND PROPERTIES

Both the chemical composition of a steel and its “thermal history” must be known in order to identify the microstructure and mechanical prop-

erties obtainable.

It

is this great flexibility in controlling the structure and

properties of steels that makes steel such a useful and versatile mate- rial. Because the thermal history of different castings may vary and even the thermal history of differing section sizes within a casting may vary, different transformations, microstructures, and mechanical prop-

erties can be expected for a given alloy. Figure 4 illustrates the effects of various heat treatments on the mechanical properties of cast steels with various carbon contents. Microstructure differences are illustrated in Figure 5 for a 0.23% C cast steel given various heat treatments. In

each case the phases present in the microstructure are the same (α +

Fe₃C), but the distribution and shape of the phases are very different re-

sulting in very different mechanical properties.

Heat treatments to produce the different microstructures in Figure 5 all

begin in the austenite region of the phase diagram with all of the carbon distributed uniformly in solid solution in the austenite phase. A high tem- perature photomicrograph of a steel heated into the austenite phase re- gion would show featureless grains of austenite with little or no evi-

dence of the prior α +Fe3C structure that was present at room

temperature. This steel is now at the "starting line" waiting for phase

transformations to occur upon cooling to transform the single phase

austenite structure to any one of a variety of room temperature micro- structures, depending on cooling conditions.

The effects of cooling rates from the austenite phase region on the

phase transformations that occur are shown schematically for the ex-

treme cases of very rapid and very slow cooling in Figure 6. Slow or

in ferrite and carbide phases forming from austenite as can be calcu- lated from the phase diagram. The microstructure obtained would con-

sist of grains of ferrite (α) and grains of pearlite (a fine two phase mix-

ture of α and Fe₃C). However rapid cooling or quenching from the

austenite phase region can prevent the formation of ferrite and carbide and can result in the formation of a new non-equilibrium phase, not on the phase diagram, namely martensite. Most are familiar with the me- chanical properties of this hard, brittle phase formed when a steel is quenched rapidly. Much is known about the physical characteristics the martensite phase and its crystallographic relationship to the other

phases of steel (4,5,6,7) but this discussion is beyond the scope of this

paper. Furthermore, if the martensite phase is given a low temperature

tempering heat treatment (in the α + Fe3C region of the phase diagram

at temperatures less than 1340°F), it will also transform forming a very

fine uniform distribution of phases αand Fe3C commonly known as tem-

pered martensite. In most cases tempered martensite has a superior

combination of strength and toughness for a given steel compared to

other α +Fe₃C structures such as pearlite. High magnification views of

pearlite formed from slow or equilibrium cooling and tempered marten- site formed from quenching and tempering are shown in Figure 7.

These details cannot be seen in conventional optical photomicrographs

as in Figure 5. Although pearlite and tempered martensite have differ- ent properties and were formed from different transformation reac- tions, both are distributions of the phases ferrite and carbide and are thus very similar. (Bainite is also an intermediate phase mixture of fer. rite and carbide that can be considered as a transition structure be-

much commercial importance to the steel casting industry it will not be discussed in any detail.)

At this point it is important to clarify terminology to note that not all quench and temper heat treatments result in the formation of tempered martensite microstructures. If cooling from the austenite phase region is not rapid enough upon quenching (for example, at the center of a thick casting) a fine structure of ferrite and pearlite will form instead of martensite. Tempering will not substantially change the structure of this fine ferrite and pearlite. Although the properties of pearlitic structures obtained from quenching and tempering will not be as desirable as for tempered martensite structures, the properties will be substantially bet- ter than for coarse ferrite and pearlite obtained from slower cooling. Heat treaters use the terms quench and temper to refer to the physical processing of the casting which may or may not produce tempered martensite. Metallurgists are more restrictive and assume by definition that cooling rates during quenching are rapid enough to form marten- site that can subsequently be tempered to form tempered martensite. It is clear that neither the phase diagram nor the simple diagram in Fig-

ure 6 is adequate to completely describe transformation behavior.

Questions such as “How slow does slow cooling have to be to get

pearlite?” or “What cooling rates are required to form martensite?”

need to be answered in quantitative ways. To answer these questions austenite transformation diagrams known as time, temperature trans- formation diagrams (TTTdiagrams) and continuous cooling transforma-

tion diagrams (CCT diagrams) need to be understood. Just as the phase

diagram is the “road map” for understanding the direction of equili-

brium transformations, TTTand CCTdiagrams are the road maps

for

un-

derstanding the time and temperature transformation response under non-equilibrium conditions that go beyond the phase diagram. These transformations will first be discussed in a classical way and then dis-

cussed in a

more

practical way as it relates to developing microstruc-

tures and mechanical properties upon heat treatment of cast steels.

AUSTENITE TRANSFORMATlON

Most heat treatments begin at high temperatures in the austenite re- gion of the phase diagram. During austenitization the initial structure of the steel, whether as-cast or already heat treated, is “erased” and re- placed by a uniform austenite solid solution. The details of the transfor- mation of initial structures to austenite are of some interest because

this transformation rate dictates the necessary time to fully austenitize a casting and dissolve all of the carbide phase. Recent SFSA research has indicated that standard austenitizing heat treatment practice may be overconservative and that shorter heat treatment cycles than have been used in the past may be adequate to fully austenitize steel cast- ings (1 1).

However; it is the transformation of austenite upon cooling that must be studied in detail to understand the wide variety of microstructures and resultant mechanical properties that can be obtained upon cooling. The phase diagram indicates that under equilibrium conditions it is not pos-

sible to have austenite present at temperatures below 727°C (1 340°F).

But when austenite is cooled to below the transformation temperature it

can take a considerable amount of time before ferrite and carbide form.

It is this time delay that allows austenite to be commonly cooled sub-

stantially below 727°C (1 340°F) and to transform at various tempera-

tures even though the phase diagram does not indicate that this is pos-

sible.

It

is the temperature or temperatures at which this transformation

takes place that determines the final microstructure.

Two basic parameters determine the speed of the transformation from undercooled austenite in steel: the nucleation rate and the growth rate

of the newly forming phase(s). The term nucleation rate refers to the

rate of initiation of a phase transformation at discrete sites in the micro- structure. In practical terms, the nucleation rate is determined by the amount of undercooling (more undercooling favors nucleation) and by the rate of diffusion of the atoms participating in the transformation. Substantial diffusion of carbon is required during transformation be-

cause austenite typically with 0.2 to 0.6% C must form ferrite (0.02%

C) and carbide (6.7% C). The growth rate of a new phase, once nuclea- tion has occurred, is determined primarily by the transformation tem- perature which determines the mobility of the atoms. Much more de- tailed descriptions of nucleation and growth concepts can be found in numerous physical metallurgy textbooks including reference (4). Characteristic isothermal transformation from austenite to ferrite and carbide at a temperature below 727°C (1340°F) is shown schemati- cally as a function of time in Figure 8a and shown in somewhat more detail in Figure 8b for the transformation of a eutectoid steel (0.77% C) at selected transformation temperatures. These transformation curves can be observed experimentally using a number of techniques includ-

ing dilatometry or electrical resistivity measurements (8). The “S”

shape of these curves is characteristic for all nucleation and growth

type reactions. Note that in Figure 8b, nucleation occurs most rapidly at the intermediate transformation temperature of 1000° F and more slowly at 800°F or 1300°F.

A much more common way of representing austenite transformation characteristics upon cooling is shown in Figure 9 where the austenite transformation response of a eutectoid steel is shown as a function of time and temperature with characteristic “C”shaped curves on a tem- perature vs. log time plot. Nucleation of ferrite and carbide from the original austenite (transformation start) occurs most rapidly at interme- diate transformation temperatures where both high nucleation driving force and relatively high atom mobility favor rapid transformation. The distribution of the ferrite and carbide phases formed isothermally from austenite at the various transformation temperatures is much dif- ferent even though the phases themselves and their relative amounts are the same, as shown at very high magnification in Figure 10. As the

transformation temperature decreases the α + Fe3C phase distribution

formed gets finer and finer and a gradual microstructural change from

coarse pearlite to fine pearlite to upper bainite to finally to lower bainite

is observed. (The terminology is complex, but the α

+

Fe3C phases

formed are the same in all cases.) As the α

+

Fe3C structure becomes

finer and finer at lower transformation temperatures the hardness and strength of the steel increase. Therefore it is not only the amount of the

Fe₃C phase present (as determined by a steel’s carbon content) but the

distribution and fineness of the Fe₃C phase (as determined by the trans-

formation conditions) that determines the strength and hardness of a steel. Figure 10 is not a complete isothermal transformation diagram (TTT diagram) for this eutectoid steel because the transformation of austenite to martensite is not shown on the diagram.

The transformation of austenite to form ferrite and carbide can be

avoided if a steel is cooled rapidly enough to suppress the nucleation

and subsequent growth of ferrite and carbide as either pearlite or bainite. The martensite phase that forms when cooling austenite very rapidly and the resultant martensitic transformation is substantially dif-

ferent from the α + Fe3C nucleation and growth transformations just

described. Martensite forms by a diffusionless shear transformation from austenite and has the same carbon content as the parent aus- tenite. The intricacies of the martensitic reaction and the properties of

this hard, brittle carbon supersaturated phase are complex but are

described elsewhere (4-8). The amount of martensite formed upon

quenching a steel that has been austenitized is a function only of the

quenching temperature. Austenite transforms instantaneously to mar-

tensite at any transformation temperature below the martensite start

temperature (Ms temperature).

A complete TTT diagram for a 0.45% carbon steel is shown in Figure 11

along with the corresponding portion of the phase diagram for this steel. An additional region of transformation from austenite t o ferrite is observed in the upper portion of the diagram. This region occurs only

for

hypoeutectoid steels and corresponds to the formation of some fer-

rite prior to the formation of pearlite at high transformation tempera- tures. If austenite is quenched rapidly enough to avoid the formation of

α + _Fe3C, it will transform to martensite as indicated in the lower por-

tion of the diagram. Quenching to the martensite finish temperature (M_f

temperature) or below will result in the transformation of all of the aus-

temperatures will result in structures with retained austenite present along with martensite. The amount of retained austenite remaining in the microstructure after quenching is of practical importance particu-

larly for alloyed high-carbon steels where the Mf temperature may be

considerably below room temperature. In general significant amounts of retained austenite are not desirable.

It

is important to use and understand these transformation diagrams

correctly TTTdiagrams indicate only the transformation of austenite to

ferrite, pearlite, bainite, martensite, or combinations of these struc- tures. Non-austenite microstructures that form do not further transform

if cooling passes through another region on the TTTdiagram. For exam-

ple pearlite, once formed from austenite, cannot transform to bainite or martensite.

In document Steel Casting Metallurgy (Page 51-59)