\input texp \nopagenumbers 



\centerline{\font\tenbf=times at 18pt\bf {MECHANICAL PROPERTIES OF
MARTENSITE}}
\centerline{\font\tenbf=times at 18pt\bf {IN HEAT--RESISTANT STEELS}}
\bigskip
\centerline{\font\tenrm=times at 14pt\rm H. K. D. H. Bhadeshia}
\centerline{\font\tenrm=times at 14pt\rm University of Cambridge, U.K.}
\bigskip \singlespace


\ssec{\bf ABSTRACT}

It is interesting that martensite is the key microstructure in
the most modern of creep resistant steels. This may be an
accidental consequence of alloy design, or it may be a
fortunate outcome.  For
the energy generation industries, the temperatures are such
that creep occurs with the climb of dislocations over
obstacles. It is argued that in these circumstances, it is
favourable to have a fine martensite plate microstructure.
Factors controlling the size of martensite plates are discussed
and novel experiments are proposed to identify the role of the
martensite plate boundaries. At the same time, there are some
detrimental consequences of a martensitic or bainitic
microstructure which need to be controlled for optimum
properties. 

\ssec{\bf INTRODUCTION}

The most exciting ferritic creep--resistant steels start with a
microstructure which is martensitic [1,2]. On the other hand,
the best established steels of this kind rely on allotriomorphic
ferrite or bainite as the starting microstructure. It is therefore
pertinent to ask: 

{\parindent=25pt \narrower \medskip
\zz{(a)} Why are modern heat--resistant steels based on
martensite?
\zz{(b)} What makes martensite  desirable for heat--resistant
steels?
\zz{(c)} Can the martensite itself be altered to better fit the
application?
\medskip}

\ssec{\bf WHY MARTENSITE?}

Low--chromium steels, such as the classical ${\rm 2{1\over
4}Cr1Mo}$ or 1CrMoV alloys have formed the backbone of
the power generation and petrochemical industries for at least
five decades, for operating temperatures of 565\deg or less.
The ${\rm 2{1\over 4}Cr1Mo}$ is essentially bainitic, whereas
the 9Cr1Mo type alloys developed much later, for higher
temperatures and greater corrosion/oxidation resistance, are
martensitic. They undoubtedly have better creep strength
(\fagg a, b), but why are they martensitic? 

The probable answer to this has little to do with
microstructure. A greater chromium concentration is needed to
obtain the oxidation and corrosion resistance necessary for
higher service temperatures. The chromium  must then be
balanced by other solutes to avoid an excessive fraction of
$\delta$--ferrite. The net solute content then becomes so
large that the steels cannot transform to bainite and hence
the martensitic microstructure. This is illustrated by the
time--temperature transformation diagrams for 2.3, 4.3 and
9.3 wt\%\ chromium steels shown in \figg c.

\picture{grand}{311}{270}{400}{\tabtit{\figg :
}{ (a,b) Calculated rupture
strength as a function of time for a bainitic (low Cr) and
martensitic (high Cr) steel. The details can be found in [3]. (c) Calculated time--temperature transformation
diagrams for steels {\bf x}
(0.15C--0.25Si--0.5Mn--1Mo--2.3Cr wt\%), {\bf y} (4.3Cr)
and {\bf z} (9.3Cr).  The transformation curves refer
to zero percent reaction. In each case the upper
curve represents diffusional transformation whereas
the lower curve represents bainite.}}

Could it be that the martensitic microstructure, which has a
very high number density of defects,  encourages the
precipitation of more numerous and finer carbide particles
which resist creep deformation? Baker and Nutting [4]
conducted a transmission electron microscope study
of carbide precipitation reactions during the isothermal
tempering of a ${\rm 2{1\over 4}Cr1Mo}$ steel in two
starting microstructural conditions, bainite and martensite.
\fagg\ shows that whereas the kinetics of precipitation are
not identical in the two cases, the differences are very small
over the temperature range of interest (500--650\degg).
This is to be expected since both microstructures contain
similar defect densities (dislocations, plate boundaries,
austenite grain boundaries); the slightly smaller defect density
expected in bainite does not seem to have any significant
effect on the precipitation kinetics. A martensitic
microstructure is not therefore necessary in heat--resistant
steels; bainite is adequate. The modern trend
towards martensitic creep--resistant steels is associated with
the need to improve the environmental resistance of the
alloys rather than microstructural considerations. Indeed, there
are well known disadvantages to high hardenability martensitic
steels when it comes to weldability.

\picture{baker}{267}{118}{400}{\tabtit{\figg : }{Measured
isothermal transformation diagrams for carbide precipitation
reactions in ${\rm 2{1\over 4}Cr1Mo}$ steel. (a) Martensitic
starting microstructure; (b) bainitic starting microstructure.
Adapted from  Baker and Nutting [4].}}

\ssec{\bf DESIREABLE FEATURES OF MARTENSITE}

Extremely fine dispersions of carbides can be achieved during
the transformation of ferrite to austenite.   \lq \lq Interphase precipitation" is a term used to describe
the precipitation of a third phase at the $\alpha/\gamma$
interface during the transformation of austenite to ferrite. The 
phase may be cementite, alloy  carbides or other particles such
as those of copper or gold which have limited solubility in iron
[5].  The scenario is one of ferrite grains on the
scale of tens of microns containing the fine dispersions of
precipitates.

It is surprising therefore that these coarse grained
microstructures have not been developed for creep resistant
applications. The martensitic or
bainitic steels  have the advantage that the plate
size is very small, of the order of a quarter of a micron. The
mean free path for dislocations is only about twice the
thickness of the plate [6]. This must provide additional
hinderance to the passage of dislocations during creep, over
and above the barriers due to precipitates. Indeed, dislocation
cell structures develop within the martensite boundaries. 
Anything that can be done to refine these plates must
therefore lead to an improvement in the creep
resistance.

{\parindent 20pt \narrower \medskip

\x \it It is proposed here that when the mechanism of creep is
one controlled by the climb of dislocations over precipitates,
there is an advantage to refining the matrix grain size.

\medskip}

\x We shall return to this proposal when considering
oxide--dispersion strengthened iron alloys. But for the
moment, what determines the thickness of martensite or
bainite plates and how can that be refined? This can be
dealt with using a quantitative model
developed for bainite [7], but similar criteria should
apply also to martensite, bearing in mind that the
transformations are both displacive in character.


\ssec{\bf THICKNESS OF PLATES}

If the shape deformation 
associated with martensitic or bainitic transformation is
elastically accommodated then the plates may be expected to
maintain an elastic equilibrium with the matrix; they may
continue to thicken until the strain energy balances the
available free energy. It follows that if the plates are allowed to
grow freely (\ie without impingement), they would be expected
to be thicker at lower temperatures, where the driving force is
the greatest. Not surprisingly, this contradicts the experimental
data  which repeatedly indicate finer plates at lower
temperatures. This is because in steels of the type considered
here the shape deformation is not elastically accommodated. 
The dislocation debris associated with the plastic
accommodation resists the advance of the bainite/austenite
interface, the resistance being greatest with strong austenite.
The strength of austenite must therefore feature in any
assessment of plate size; indeed, this effect has been known
for some time [8]. The plates are expected to become
thicker at high temperatures, where the austenite has a lower
yield strength. Dynamic recovery  at
high temperatures should further weaken the austenite and
lead to coarser plates.

The thickness must also be influenced by impingement
between adjacent plates; a large nucleation rate must
naturally refine the microstructure.

The observed effect of temperature 
could easily be indirect since both  the austenite strength and
the nucleation rates are strongly dependent on temperature
and the composition of the steel. A quantitative analysis of the
bainite plate thickness as a function of the strength of the
austenite, the chemical free energy change accompanying
transformation and the transformation temperature has shown
that the latter has only a  small independent effect on the
thickness (\fagg). Strong austenite and high
driving forces lead to a finer microstructure.

\picture{significance}  {297} {245} {450}
{\tabtit{\figg : }{(a) The model perceived significance
of each of the variables plotted on the horizontal axis, in 
influencing the thickness of bainite plates. The vertical scale
represents the ability of the variable to explain variations in
plate thickness. (b) Variation in thickness with the chemical
driving force. (c) Variation in thickness with the strength of
the austenite. After Singh and Bhadeshia [7].}}

Controlled experiments are needed to discover whether the
creep resistance of steels increases as the plate thickness is
reduced.  Obviously, any refinement of the microstructure
should not be so large that the mechanism of creep changes,
from one controlled by the climb of dislocations over obstacles,
to one dominated by diffusion along martensite boundaries.
However, this is unlikely to be an issue because the boundaries
between martensite plates in a given austenite grain tend to
have a high degree of coherence.

This naturally brings us to a consideration of the crystallography
of martensite. Davies and Magee in 1970 [9] reported some
fascinating experiments in which the {\it crystallography} of
martensite was studied as a function of the magnetic state of
the austenite, and when the austenite was solid--solution
strengthened using either substitutional or interstitial solutes. I
am not aware of any work relating martensite
crystallography to the creep properties of steels.

\ssec{\bf OXIDE DISPERSION STRENGTHENED MARTENSITE}

Mechanical alloying  [10--13] is a process in which an alloy is
created  by the intense
mechanical deformation of mixtures of elemental or
master--alloy powders (\fagg). The powders are then
consolidated by various combinations of hot--isostatic pressing
and extrusion; the alloys in this condition  are usually too hard
to use. They are therefore heat--treated, either isothermally or
in a temperature gradient to induce recrystallisation. Mechanical
alloying permits  the usual limits of solubility imposed by
solidification to be exceeded, but in addition, dispersoid
particles (such as oxides) can be introduced uniformly into bulk
materials. 

\picture{grand2}{290}{266}{410}{\tabtit{\figg : }{(a) A
schematic illustration of a common method of manufacturing
large quantities of mechanically alloyed metals for engineering
applications. The elemental powders/master--alloys/oxides are
milled together to produce solid solutions with oxide
dispersoids. This powder is then consolidated and the resulting
bulk alloy heat--treated to achieve a coarse, directional grain
structure. (b) An ODS alloy with coarse matrix grains. (b) A
martensitic ODS alloy.}}

Mechanically alloyed iron--base materials are used in a
recrystallised condition with extremely coarse grain structures
[14]. This is because the are destined to operate in environments
where the temperature is very high, about 1000\degg, where
creep by diffusion along boundaries is an important
contribution to the overall creep strain. 

Typical commercial alloys are ferritic at all temperatures. A
critical esperiment to test the role of martensite
boundaries in creep resistant steels could involve the
manufacture of heat treatable ODS alloys [15]. This would
allow the testing of a dispersion--strengthened martensite
(\figg b,c) which can  be compared using low temperature
(600\degg) creep measurements in two states, the first  a
recrystallised grain structure and the second with an identical
oxide dispersion but in a martensitic matrix. The experiment
would reveal directly the role of the martensite boundaries in a
creep process controlled by dislocation climb over obstacles.

\ssec{\bf UNDESIREABLE FEATURES OF MARTENSITE}

\ssec{\bf Grain Boundary Embrittlement}

During diffusional growth, grains
of ferrite grow freely across the austenite grain 
boundaries, which are destroyed in the process.  By
contrast, displacive  transformations occur with 
the disciplined motion of atoms. Such movements  cannot be
sustained across grain boundaries. It follows that a vestige  of
the austenite grain boundary remains when the transformation
is displacive. These prior boundaries become susceptible to
impurity segregation and embrittlement, liquid  zinc
embrittlement, and any phenomenon able to exploit the
disorder associated with these boundaries [16--19]. 

The disorder at the prior austenite grain boundaries is larger
than that associated with martensite--martensite boundaries.
This is because adjacent martensite plates have coherence
with the parent austenite grain and hence a good fit with each
other, whereas the orientation relationships between
austenite grains can be random. There are considerable
practical consequences of this difference. Thus, the extent of
segregation is smaller at martensite lath boundaries. Carbides
nucleate preferentially at the prior austenite grain boundaries
during the tempering of martensite or bainite [20,21].  The
boundaries can be revealed by etching
[22], often with great clarity even though the original grain is
no longer present.  

The problem of austenite grain boundary embrittlement in
martensitic or bainitic steels is not exclusive to high strength
steels used at ambient temperature.  It is well established that
the fracture toughness of many power plant steels deteriorates
during service at elevated temperatures for two reasons.
Firstly, the carbide particles, particularly those located at
at the prior boundaries, coarsen and hence provide
easier sites for crack or void nucleation. Secondly, the
elevated temperatures permit the impurities to
diffuse relatively rapidly and saturate the boundaries.  \fagg\
shows the impact transition curves for both the virgin steel
(\ie at implementation into service) and for an ex--service
sample which had experienced 88,000~h at 813~K [23]. 

\picture{impact}{137}{114}{500}{\tabtit{\figg : }{Impact
test data on 2${1\over4}$Cr1Mo steel. (a) Comparison of the
impact energies of the steel in its virgin condition, after
service exposure at 813~K for 88,000~h and for a
service--simulated sample (873~K for 10,000~h). After Wignarajah
\et [23].}}




\ssec{\bf Avoiding Grain Boundary Embrittlement}

The traditional routes for avioiding temper
embrittlement involve a reduction in the impurity concentration
and the addition of substitutional solutes which getter the
impurities [24].

A solute causes embrittlement by reducing the cohesion
across the boundary plane. On the other hand, those solutes
which enhance covalent bonding normal to the boundary must
reduce embrittlement. {\it Ab--initio} calculations of the
bonding at grain boundaries in iron suggest this mechanism and
confirm  that phosphorus should embrittle and boron should
toughen [25].

 Boron additions must therefore be useful in
commercial, martensitic power plant steels, which can be made
clean but not free of impurities. The difficulty, however, is that
much of the boron  enters the ${\rm M_{23}C_6}$ phase
[26] and  hence less is available to segregate to the
austenite grain boundaries. Much further work is needed on
this topic.

Another possibility, which has never been investigated, would
be to reduce the energy of the prior austenite grain
boundaries. This could be done by making  the austenite
grain structure strongly crystallographically textured. The
resulting austenite grain boundaries would have a high degree
of fit and hence  be less prone to impurity segregation.

\ssec{\bf DOUBLE AGEING HEAT TREATMENTS}

There is a tendency during the tempering of martensite for precipitation to
occur preferentially at the prior austenite grain boundaries and at the
martensite plate boundaries, because these represent defects 
where the activation energy for nucleation is smallest. A given
material will contain nucleation sites with different activation
energies. However, the relative nucleation rates do not necessarily increase
in the same order as the activation energies because the number density of
high activation--energy sites may be much larger [27].
Hence, more uniform microstructures are possible if heat
treatments are conducted at large driving forces, which in the
context of martensitic steels corresponds to lower
temperatures.

The important precipitates in heat--resistant martensitic
steels are rich in substitutional solutes and hence cannot grow
at low temperatures. Furthermore, many heat--treatments
have the dual purpose of relieving residual stresses and
undesirable local microstructures arising during fabrication
operations such as welding. The minimum temperature at which
substitutional atoms in iron are mobile over reasonable time
scales is around 500\deg (\fagg a), whereas commercial heat
treatments are conducted at temperatures in excess of
700\degg. It may be possible to produce a better dispersion of
precipitates by employing a dual heat treatment procedure in
which the initial anneal is at 500\deg followed by an elevated
temperature treatment at 720\degg.  

Heat--treatment at the lower temperature (larger driving
force) should lead to more uniform precipitation and a
larger number density of precipitates, whereas the
subsequent higher temperature treatment allows the
precipitates to grow, and for appropriate diffusion to anneal
residual stresses and brittle microstructures such as those
induced in the heat--affected zone of a weld.

\picture{diffuse}{256}{99}{500}{\tabtit{\figg : }{(a) The random walk
diffusion distance (in nanometers) of elements in ferritic or austenitic iron, at
500$\,^\circ$C and 300$\,^\circ$C. (b) Dual heat treatment.}}

On the other hand, the situation may not be as clear as
implied here; research on plain carbon steels indicates that
particles at lath boundaries stabilise the lath microstructure.
The microstructures of both tempered martensite and bainite
contain two kinds of cementite particles, those located at the
lath boundaries and a finer distribution within the laths. In
upper bainite the cementite is only located at the lath
boundaries. \fagg\ shows experimental data on the coarsening
of cementite during the tempering of a medium carbon steel.
The upper bound of each shaded region represents the
lath--boundary cementite, the lower bound the intra--lath
cementite. The bainitic microstructure is coarse to begin with
because of the tempering inherent in the formation of bainite.
With martensite the tempering induces the precipitation of
cementite, with considerable intra--lath cementite and a larger
overall number density of particles. Therefore, the coarsening
rate is much larger for martensite; the bainitic starting
microstructure shows considerable stability to tempering. A
consequence is that the martix microstructure remains refined
over a longer time period for bainite than for martensite.

\picture{cementite}{154}{106}{500}{\tabtit{\figg : }{Changes
in the size of cementite particles as a function of the
tempering time at 700$\,^\circ$C, with different starting
microstructures. The upper bound of each shaded region
represents the mean size of particles located at lath
boundaries. The lower bound corresponds to particles within
the laths. The data are for a Fe--0.45C--0.22Si--0.62Mn wt\%\
steel; the bainite was prduced by isothermal transformation at
380$\,^\circ$C. After Nam [28].}}


\vfill\eject \ssec{\bf CONCLUSIONS}

The fact that modern creep resistant steels have a
predominantly martensitic microstructure may be a
fortuitous consequence of efforts to improve the corrosion
resistance. It is nevertheless likely that such a microstructure
had advantages, for example, the plate boundaries may 
provide additional barriers to dislocation creep. However,
this hypothesis is not proven so some interesting experiments
are suggested as critical tests. It is only when the hypothesis
is established would it be reasonable to attempt
modifications to the size, shape and crystallography of
martensite as a means towards better creep resistance. 

The high sensitivity of martensitic microstructures to impurity
embrittlement can be controlled by minimising the deleterious
elements in the steel. However, alternative approaches, such
as the modification of the crystallographic texture of
austenite in order to improve the fit at the prior austenite
grain boundaries, could prove worthwhile.

\ssec{\bf ACKNOWLEDGEMENTS}

I am  grateful to the organisers of this workshop for the
opportunity to present this work. I would also like to thank
them for the  generous financial support. I would never have
thought about the specific role of   martensite in creep
resistant steels without Dr Kaneaki Tsuzaki's suggestion. 

\ssec{\bf REFERENCES}

{\parindent=10pt \narrower \medskip
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\medskip}
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