(3 13 favourable and expected to form more readily on dislocations.

Also, precipitation of TiC can occur at stacking faults,,

2.h*5 Precipitation of 'f1

In general, if the Ni content of alloys is of the order of 15-3Q?o and with the normal A1 and Ti additions of upto

k% of each element, the Ni/(A1+Ti) ratio is sufficiently high

( 39)

f or Vjr* precipitation to occur. J J Although the lattice

•parameter of the matrix changes slightly with small changes in composition, the difference in lattice parameter between T

and matrix is very slight, generally not more than 0. .

v/ith the matrix, and the normal grov/th mechanisms operate. Generally, in the case of a completely coherent precipitate and smaller particles*the process of coarsening has been

(77)

reported to be diffusion controlled.v / Although for

semi-coherent interfaces, a simple ledge-growth mechanism has been proposed^^ , it is argued that it may not necessarily

(77)

be operative for coherent interfaces. '

Usually X has very small matrix misfits, but however,

small this difference may seem, it creates considerable internal strain during the period of coherency, when the lattice of the matrix is forced to conform to the lattice of the precipitate.^-^' The coherent particles have been reported to have symmetrical hexagonal dislocation networks at the interface, v/hich along Y/ith the particle dimensions; have been regarded as the major factor controlling the loss of coherency between the particle and the matrix.^79) Although there have been three different mechanisms suggested for the loss of coherency of a growing

particle 9 . in the case of y\ ^as heen shown that coherency

is lost by the absorption of matrix dislocations at its inter­

face. (79) operative mechanism in the case of <]ff has been

described by the interaction between the strain field of the dislocation and the elastic strain field of the precipitate. The precipitate is effectively metastable with respect to the dislocation if r is > rcritical wiiere*“

rcritical = ra(^ius ^ e particle belov; v/hich a

precipitate cannot support a dislocation loop at its interface.

and r = actual radius of the particle at a particular

stage of ageing.

The shape of the y % precipitate depends largely upon the

matrix/particle misfit. The shape can be either spherical

or cubic and in most cases when the misfit between ^ct and austenite is quite small, a spherical precipitate is formed, which in general shov/s a very slow rate of grov/th, and has a cube-cube orientation relationship with the matrix.

Systems with misfits around 0*5% are at the critical stage

The shape is also controlled by the size of the particle, although the misfit can be altered by small changes in matrix

solute during ageing. Particle dimensions can also contribute

to this parameter to a large extent, thereby showing the ( 77)

importance of particle size.x,,/ 2.h«»5ol Precipitation Sites for y*

It is widely known that nucleation of solid state

transformations occurs on lattice defects due to the release of the excess free energy as well as reducing the total

interface energy required^^. Some of the accepted sites of precipitation for 1 are:-^-^

a) Intragranular precipitation and b) Precipitation on dislocations

Homogenous nucleation of <jrf takes place in the absence of any pre-existing defect in the matrix, as reported by Saito and W a t a n a b e . O n further ageing, ordered zones were produced by the emission of vacancies. Wilson and Pickering^^ examined the early stages of ageing of an austenitic steel which precipitated y* and reported the

hetrogeneous nucleation of y 1 on pre-existing lattice defects. Recent work^2^ on PEl6 reports the quenched in-vacancies

as the controlling factor for y* nucleation, and also shews that the y* particle, is a vacancy source in the alloy. An increased vacancy emission was observed at the particle - matrix interface, which could increase the rate of nucleation for further precipitation.

Precipitation of y*' occurs on dislocations^*^ and nucleation has been considered for both the coherent and incoherent case^2^. In the case of incoherent nuclei an equilibrium phase will be formed, and in the case of coherent nuclei a transition phase may be the result. Nucleation may, however, be suppressed near dislocations if the precipitate has little or no misfit strain and has a higher shear modulus

than the matrix, because a hard precipitate increases the strain energy associated with a dislocation. Under certain

conditions,nucleation can occur preferentially on dislocation of the edge type.^^ Dollins^-^ recently presented a treat­ ment for nuclei forming not on a dislocation line but in its vicinity, due to the catalytic effect exerted by the dislocation core.

2o!u5«>2 Effect of A1 and Ti Content on Precipitation of of*

in Austenite ~,

The effect of these alloying additions on y* precipitation

can be summarised as follows:- a) Effect of Ti

In austenitic steels containing Ti only, the rate of

overageing is rapid due to the formation of cellular precipitate

of Ni^ Ti when the Ti content exceeds 3%. ^3) jn Ni„-base

superalloys, the compound Ni^Ti,usually designated eta, has a h.c.p structure in contrast to the f.c.c. of the austenite matrix. This precipitate is f.c.c. during the period of

coherency with the f.c.c. matrix, but precipitates as particles of second phase which have either the h.c.p. (normal, lowest energy) or f.c.c. abnormal s t r u c t u r e . A s precipitation

continues the f.c.c. y x form of Ni^Ti phase dissolves and is

replaced by the h.c.p. form. Ageing at lower temperatures allows more of the f.c.c. to be precipitated than does ageing at higher temperatures. Ni^ Ti is considered to be undesirable because it has no solubility for other elements and adheres closely to the stoichiometric compositipn.

b) Effect of A1

It has been observed that a relatively small amount of age- hardening is produced by A1 alone in both austenitic steelsv 1 and Ni-base superalloys. . In austenitic steels containing 2 % Ni-15Cr and more than 1% Al, the age-hardening response

increased, but not to a great extent. The effects of precipitatic were complex, and there was evidence that <yf was f ormed.

In Ni-base superalloys the lower intensity of age-hardening

precipitation containing only Al0 However, there is evidence

that increasing A1 in austenitic steels increases the tensile

strength and 0.2$ P.Sc but decreases ductility and impact c) Effect of^Al + Ti) added together

In Ni-base superalloys, the presence of both(Al + Ti)_{/ Or\ '}

increases the hardness considerably.' The qf* phase is

capable of taking a considerable amount of other elements

into solution, namely Or , Ti and Co. Grant et. al' J have

shown that up to 65% of the A1 atoms in Ni^ A1 can be replaced

by Ti. Many details are uncertain, but is generally design­

ated as Ni^ (Al.Ti). Investigations have also revealed solubility for other elements, which results in a greater

volume fraction of ff* than could be possibly formed if the

composition were simply Ni^ (Al.Ti).

In austenitic steels containing 25%> Ni-1 5 Or, the amount

of ageing was greater in(Al + Ti)steels than could be obtained

with a single addition of similar amounts of A1 or T i . ^ ^

With low A1 contents the presence of Ti changed the precipitate

from Ni A1 to Ni^ (Al.Ti), so the Ti caused b0c.c. precipitate

to be replaced by the f.c.c. precipitate. It is well known that the Al/Ti ratio plays a dominant role in the type of precipitate formed and its orientation relationship with the

matrix.(^3) jn austenitic steels having Al/Ti ratios^ 2.6,

the over-aged precipitate was b.c.c. Ni(Al.Ti). However, Hughes found that on decreasing this ratio to about unity the

overaged precipitate was Ni^Al Ti and with a further decrease in Al/Ti ratio, Ni^ (Al.Ti) was found to be the over-aged

precipitate but no Ni^ A1 Ti was observed.

In a report on a series of alloys having different Al/Ti

ratios, and base composition of nimonic P E l6 , ^ ^ an Al/Ti

ratio of unity was found to produce the best age-hardening

characteristics. This is because the misfit between f* and

the austenite matrix increases with increased Ti (which

replaces A1 in the qf*) because of the larger atomic size of Ti

Ni^ (Al.Ti), compared with Ni^Ti, and because of its slightly

lower misfit with the matrix the rate of grov/th of is

very slow and overageing is also slow. This is considered

( 30 39)

desirable for high temperature applications® 9

2ckoS Combined Precipitation of Cg and ffx

Different heat -treatment sequences have been developed to precipitate both M23 Cg and ff1 in the same matrix, to obtain the beneficial effects of both p h a s e s . I n both austenitic steels and nimonic alloys a denuded zone has been observed, between the carbides at the grain-boundaries and in the matrix*(30,37*69) ^w0 mechanisms have been proposed for the formation of precipitate free regions, namely: a) solute

depletionfb) vacancy depletion® In the solute depletion

mechanism, it is proposed that -under ageing conditions in

which the rate of nucleation and thickening of the Cg

grainboundary precipitate is much greater than that of the intragranular <yf, the growth of the grain-boundary precipitate depletes the adjacent regions of solute, relieving solute

supersaturation, and thereby causing in this region to

d i s s o l v e ® j n vacancy depletion mechanism, it is

postulated that in the region adjacent to the grain-boundaries, during the quench from solution treatment, excess vacancies migrate to the grain-boundaries, which serve as effective

sinks for vacancies® The resultant diminution in

vacancy super-saturation effectively prevents, in these regions, the formation of vacancy loops and helical dislocations which often serve as the heterogeneous nucleation sites required for intragranular precipitation of y*» Thus, by this vacancy depletion mechanism, a precipitate free but solute rich region is produced in the vicinity of the grain-boundary® Hov/ever, these processes can be competitive and both occur simultaneously during the ageing of an a l l o y ® T h e r e f o r e , depending upon the activation energy for solute diffusion and vacancy migration, it should be possible to predict the mechanism which dominates in a particular system,for a particular heat-treatment®