Crystallography of martensite formation Martensite, the hardening constituent in quenched

8.3.3.1 Crystallography of martensite formation Martensite, the hardening constituent in quenched

steels, is formed at temperatures below about 200 ° C. The regions of the austenite which have transformed to martensite are lenticular in shape and may easily be recognized by etching, or from the distortion they pro- duce on the polished surface of the alloy. These relief effects, shown schematically in Figure 8.21, indicate that the martensite needles have been formed not with the aid of atomic diffusion but by a shear process, since if atomic mobility were allowed the large strain energy associated with the transformed volume would then be largely avoided. The lenticular shape of a martensite needle is a direct consequence of the stresses produced in the surrounding matrix by the shear mechanism of the transformation and is exactly analogous to the sim- ilar effect found in mechanical twinning. The strain energy associated with martensite is tolerated because the growth of such sheared regions does not depend on diffusion, and since the regions are coherent with the matrix they are able to spread at great speed through the crystal. The large free energy change associated with the rapid formation of the new phase outweighs the strain energy, so that there is a net lowering of free energy.

Direct TEM observations of martensite plates have shown that there are two main types of martensite, one with a twinned structure (see Figure 8.22), known as

Figure 8.21 Schematic diagram of the observed shape deformation produced by a martensite plate. (a) Contour lines on an originally flat surface; (b) section of the surface through AB. Vertical scale much exaggerated (after Bilby and Christian, 1956; courtesy of the Institute of Metals) .

acicular martensite, and the other with a high density of dislocations but few or no twins, called massive martensite.

In contrast to the pearlite transformation, which involves both a redistribution of carbon atoms and

a structural change, the martensite transformation involves only a change in crystal structure. The structure cell of martensite is body-centred tetragonal, which is a distorted form of a body-centred cubic structure, and hence may be regarded as a super- saturated solution of carbon in ˛-iron. X-ray examina- tion shows that while the c/a ratio of the bct structure of martensite increases with increasing carbon content, the curve of c/a ratio against composition extrapolates back to c/a D 1 for zero carbon content, and the lattice parameter is equal to that of pure ˛-iron (Figure 8.23).

From the crystallographic point of view the most important experimental data in any martensite transfor- mation are the orientation relations of the two phases and the habit plane. In steel, there are three groups of orientations often quoted; those due to Kurdjumov and Sachs, Nishiyama, and Greninger and Troiano, respec- tively. According to the Kurdjumov–Sachs relation, in iron–carbon alloys with 0.5–1.4% carbon, a f1 1 1g plane of the austenite lattice is parallel to the f1 1 0g ˛

plane of the martensite, with a h1 1 0i axis of the former parallel to a h1 1 1i ˛ axis of the latter; the asso- ciated habit plane is f2 2 5g . In any one crystal there are 24 possible variants of the Kurdjumov–Sachs rela- tionship, consisting of 12 twin pairs, both orientations

Strengthening and toughening 279

Figure 8.22 (a) Formation of a martensite platelet in a crystal of austenite; (b) the inhomogeneous twinning shear within the platelet (after Kelly and Nutting, 1960; courtesy of the Royal Society) .

of a pair having the same habit plane. However, for

8.3.3.2 Mechanism of martensite formation general discussion it is usual to choose one relation

The martensite transformation is diffusionless, and which may be written

therefore martensite forms without any interchange in ⊲

the position of neighbouring atoms. Accordingly, the

observed orientation relationships are a direct conse- In the composition range 1.5–1.8% carbon the habit

1 1 1⊳ 1 0 1⊳ ˛ with [1 1 0] [1 1 1] ˛

quence of the atom movements that occur during the

transformation. The first suggestion of a possible trans- tionship unspecified. This latter type of habit plane has

plane changes to ³ f2 5 9g with the orientation rela-

formation mechanism was made by Bain in 1934. He also been reported by Nishiyama for iron–nickel alloys

suggested that since austenite may be regarded as a p (27–34% nickel) for which the orientation relationship

2, the is of the form

body-centred tetragonal structure of axial ratio

transformation merely involves a compression of the

c -axis of the austenite unit cell and expansion of the

⊲ 1 1 1⊳ k⊲1 0 1⊳ ˛ with [1 2 1] k[1 0 1] ˛

˛ -axis. The interstitially dissolved carbon atoms pre- vent the axial ratio from going completely to unity,

However, Greninger and Troiano have shown by pre- and, depending on composition, the c/a ratio will be cision orientation determinations that irrational rela-

between 1.08 and 1.0. Clearly, such a mechanism can tionships are very probable, and that in a ternary

only give rise to three martensite orientations whereas, iron–nickel –carbon alloy (0.8% carbon, 22% nickel),

in practice, 24 result. To account for this, Kurdju- ⊲ 1 1 1⊳ is approximately 1 ° from ⊲1 0 1⊳ ˛ with [1 2 1] mov and Sachs proposed that the transformation takes

approximately 2 ° from [1 0 1] ˛ , and is associated with place not by one shear process but by a sequence

a habit plane about 5 ° from ⊲2 5 9⊳. of two shears (Figure 8.24), first along the elements ⊲ 1 1 1⊳ h1 1 2i , and then a minor shear along the ele- ments ⊲1 1 2⊳ ˛ h1 1 1i ˛ ; these elements are the twinning elements of the fcc and bcc lattice, respectively. This mechanism predicts the correct orientation relations, but not the correct habit characteristics or relief effects. Accordingly, Greninger and Troiano in 1941 proposed

a different two-stage transformation, consisting of an initial shear on the irrational habit plane which pro- duces the relief effects, together with a second shear along the twinning elements of the martensite lattice. If slight adjustments in spacing are then allowed, the mechanism can account for the relief effects, habit plane, the orientation relationship and the change of structure.

Further additions to these theories have been made in an effort to produce the ideal general theory Figure 8.23 Variation of c and a parameters with carbon

of the crystallography of martensite transformation. content in martensite (after Kurdjumov, 1948) .

Bowles, for example, replaces the first shear of the

280 Modern Physical Metallurgy and Materials Engineering homogeneous strain does not do this, so that some

reasonable additional type of strain has to be added. This shear can occur either by twinning or by slip, the mode prevailing depending on the composition and cooling rate. Between carbon contents of 0.2% and 0.5% the martensite changes from dislocated marten- site arranged in thin lathes or needles to twinned aci- cular martensite arranged in plates. In the martensite formed at low C contents (e.g. Fe–Ni alloys) the thin

lathes lie parallel to each other, with a f1 1 1g habit, to form pockets of massive martensite with jagged boundaries due to the impingement of other nearby pockets of lathes. The inhomogeneous shear produced by deformation twinning occurs on f1 1 2g planes in the martensite, so that each martensite plate is made up of parallel twin plates of thickness 2–50 nm. By oper- ation of such a complex transformation mode with a high index habit plan the system maintains an invariant interfacial plane.

Because of the shears involved and the speed of the transformation it is attractive to consider that dislocations play an important role in martensite for- mation. Some insight into the basic dislocation mech- anisms has been obtained by in situ observations

Figure 8.24 Shear mechanisms of Kurdjumov and Sachs. during either cooling below M s or by straining, but

0 (b) body-centred tetragonal martensite (˛ unfortunately only for Ni –Cr austenitic steels with ), (c) cubic ferrite low stacking-fault energy (i.e. ³ 20 mJ/m (˛) (after Bowles and Barrett, 1952; courtesy of Pergamon 2 ). For

(a) Face-centred austenite with f1 1 1 g in horizontal plane,

Press) . these alloys it has been found that stacking faults

are formed either by emitting partial dislocations with

b D a/ 6h1 1 2i from grain boundaries or by the dis- Greninger–Troiano mechanism by the general type of

sociation of unit dislocations with b D a/2h1 1 0i. In homogeneous deformation in which the habit plane

regions of the grain where on cooling or deformation remains invariant, i.e. all directions in this plane are

a high density of stacking faults developed, the corre- unrotated and unchanged in length. However, in all

sponding diffraction pattern revealed cph ε-martensite. such cases the problem resolves itself into one of

On subsequent deformation or cooling, regions of ε- determining whether a homogeneous strain can trans-

martensite transform rapidly into bcc ˛-martensite, form the -lattice into the ˛-lattice, while preserv-

and indeed, the only way in which ˛-martensite was ing coherency at the boundary between them. The

observed to form was from an ε nucleus.

Figure 8.25 Electron micrographs showing (a) contrast from overlapping faults on ⊲1 1 1 ⊳; A is extrinsic and B is intrinsic in nature. (b) Residual contrast arising from a supplementary displacement across the faults which is intrinsic in nature for both faults A and B (after Brooks, Loretto and Smallman, 1979) .

Strengthening and toughening 281 Because straining or cooling can be interrupted

during the in situ experiments it was possible to carry out a detailed analysis of the defect structure formed prior to a region becoming recognizably (from diffraction patterns) martensitic. In this way it has been shown that the interplanar spacing across the individual stacking faults in the austenite decreased to the ⊲0 0 0 1⊳ spacing appropriate to ε -martensite. Figure 8.25 shows micrographs which reveal this change of spacing; no contrast is expected in Figure 8.25b if the faulted f1 1 1g planes remained at the fcc spacing, since the condition of invisibility g.R D n is obeyed. The residual contrast observed arises from the supplementary displacement R across the faults which, from the white outer fringe, is

Figure 8.27 A shear of a/6 h1 1 2 i moves atoms in the fcc positive (intrinsic) in nature for both faults and ³

structure from A sites to B sites, and after half this shear the 2% of the f1 1 1g spacing. The formation of regions structure has pseudo-bcc packing .

of ˛ from ε could also be followed although in this case the speed of the transformation precluded