5° 2.1 Low Carbon Base Composition (Alloy No.: l)'
This alloy shows no age hardening, Pig. 31, although the
softening curves previously discussed in section h.2 .1
indicate some precipitation in the temperature range 550°-750°C, which is confirmed by electron microscopical evidence. Due
to the very low carbon content in this alloy, growth of
M23C5 precipitates did not occur and neither did the precipit
ation occur to any significant extent on grain boundaries or twin boundaries.
Hov/ever, it seems that even at this low carbon content, there was still some supersaturation with respect to
and this supports the evidence of the undissolved Mg^C^ in
the 0.10$C alloy at 1100°C, which indicates that the high
nickel content has decreased the solubility of carbon in austenite. The overall decrease in hardness at lower ageing
temperatures was due to the annealing-out of quenched in
dislocations, Plate l(c) and the precipitation of Mg^Cg was
not enough to compensate for this. At higher ageing tempera tures grain growth also produced some softening.
5*2.2 Medium Carbon Bas'e Composition (Alloy 2.)
Age-hardening was clearly observed due to the increased carbon content of 0.05$, Pig. 3 2. An analysis of the tempera ture dependence of the time for achieving a given hardness leve during ageing gave an Arrhenius relationship in which the
activation energy was IhO KJ/mol (3 3 k cal/mol). This
activation energy was very similar to that for carbon diffusion in austenite, but it is strange that diffusion of carbon
should be the rate controlling mechanism because it would be expected that substitutional chromium atoms would also need to diffuse. Such chromium diffusion would require an
activation energy at least twice that o b s e r v e d ^ . However,
( & ) it is probable that more ready diffusion along grain boundaries lowered the activation energy.
The sequence of precipitation was in agreement with
( PI fiQ 7Ji}
that observed by other workers * ^ ^23^6 PreciPi‘tia'tiion
at grain boundaries was probably on dislocations lying in the grain boundary plane. This was confirmed by the fact that
the M23°6 precipitates had cube-on-cube orientation relationships
with the matrix which was in agreement with Adamson’s w o r k ^ ^ . Although, electron metallography revealed the nucleation of new dislocations at the interface between the matrix and grain/twin
boundary M23^5 precipitates, no further precipitation of any
significance occurred on such defects. The reason is thought to be associated with the low carbon content of this al^.oy
compared to those used by other workers^"^. However, on ageing at lower temperatures the grain boundaryMg^C^ was more closely spaced than after ageing at higher temperatures, due to a
greater supersaturation, more copious nucleation and less particle growth.
5.2.3 High Carbon Base Composition (Alloy 3) -
Increased age-hardening compared with alloy 2, Pig. 33 >
resulted in an increased activation energy of the order of
166 KJ/mol (i+O k cal/mol) for precipitation. This was
slightly higher than that required for chromium diffusion along
grain boundaries. ^ The initially occurred on
grain boundaries and then within the grains, Plates 2(e) and
(f). Thus bulk diffusion of chromium was necessary and the combined grain boundary and bulk diffusion of chromium may explain the observed activation energy.
Apart from more closely spaced grain boundary ^23^6
precipitates at lower ageing temperatures, Plate 3(a), the
orientation relationship of these precipitates in the high carbon alloy was different. ^ ' This would largely depend upon
the initial orientation of the dislocation lying in the grain- boundary plane. As long as the carbide has a parallel
orientation relationship with the parent austenite, precipitation
(
21)
is generally energetically favourable.' y This alloy also
within the grains, Plate 3(b), with an orientation relationship similar to that reported by other workers. "^^.The
mechanism of Widmanstatten type M f o r m a t i o n is believed
() ft)
to be the same as that proposed by Singhal and Martin, '
because these precipitates were mainly in the grain interiors rather than near the grain boundaries. Although, some
Widmanstatten ^23^6 P^cipi^ation was observed near the non
coherent edge of the twins, Plates 3(a) and (b), these are
thought to be less efficient vacancy sinks than grain boundaries
The stringer type carbides must be associated with the
dislocations'1 for such precipitates were also observed
nearer twin boundary precipitates which are known to generate dislocations during growth, Plate 3(d).
The morphology and orientation relationship of ^23C6
precipitated at non-coherent twin boundaries, Plates 3(c) and
(d), were similar to that observed by Lewis and Hattersley. ^^0
However, the coherent twin boundary ^23^6 Preo^ P ^ a^^on was
observed after ageing at longer times at higher ageing
temperatures, and this observation is in agreement with that of Adamson, Plate 3 ( ® ) * ^ ^ However, the growth of such
precipitates was slow and occurred parallel to the twin plane, rather than fanning out into the matrix. This is possible
because coherent twin JLnterfaces are poorer vacancy sinks
than grain or non-coherent twin boundaries. Nevertheless, at high ageing temperatures,^the lesa precipitation can be attributed to a smaller supersaturation of solute, and.:, to the rapid annealing out of the defects which act as nucleation sites.
5.2.4 Opmnarison of the Effect of Carbon Content -
The ageing curves, Pigs. 32 and 33 9 indicated that even
in these low carbon steels a two stage ageing process was operative. The microstructures of both the medium and high carbon alloys showed that precipitation first occurred on the grain boundaries but later at higher ageing temperature or longer ageing times, spreads into the matrix. This is thought to be the cause of a two stage age-hardening process.
3c 2.3 Bff'ect of Molybdenum (Alloy 4)
"carbides considerably lowered the amount of carbon available for precipitation, and thus the intensity of age hardening, Table V. This agrees with the microstructures which showed
only grain boundary precipitation and very little matrix
precipitation, Plates 4(b), (c) and (e). It is also significant
that the age-hardening did not show a two stage process. The occurrence of a rapid reaction at lower ageing temperatures of 500-600°C but a slower reaction at higher temperatures has already been discussed in terms of the annealing out of the dislocations generated by the undissolved carbides, section 4«2.4. This is further supported by electron microscopy,
Plate 4(e), in that at low temperatures precipitation occurred on dislocations but at higher temperatures the dislocations annealed out before precipitation occurred and so precipitatio was mainly on grain boundaries, Plate 4(c). The temperature dependence of these two reactions supports this suggestion in that the high temperature reaction had an activation energy
of 106 KJ/mol (25 K cals/mol) typical of grain boundary
m23°6 precipitation in the molybdenum - free alloy. On the
other hand, the low temperature reaction had a much higher
activation energy of 720 KJ/mol (171 K cals/mol), which
is high even for precipitation on dislocations which would require bulk solute diffusion.
A tentative *0-curve1 for the start of carbide precipitation is shown in comparison with the molybdenum free alloys in
Big. 39. The rate of the reaction seems to be about the same as for the same carbon content in the molybdenum free steel,
but the nose of the steel has been depressed by about 100°G.
It is possible that the C-curve for the molybdenum steels
is a combination of the two ageing reactions already discussed, but the precise mechanism for the depression in the nose of
the 0-curve is at present very uncertain.
However, on comparing Plates 3(d), and 4(d), it is
evident that the. molybdenum bearing steel‘.contained many fewer but coarser precipitates along non-coherent twin boundaries.
This is perhaps a direct reflection of the amount of carbon available for precipitate nucleation and growth.
5.2.6 Effect of Titanium (Alloy Nos. 5 and 6)
The two alloys which contained 1.2#Ti with^3.5$Mo,
contained very different carbon contents of 0.11$ (Alloy 5)
and 0.35/b (Alloy 6). The presence of the higher carbon
content resulted in a greater volume fraction of undissolved TiC, Table IX, and whilst in the low carbon steel there may be just sufficient titanium available for a small amount of ^precipitation, this is certainly not the case in the
/
higher carbon alloy. N o f precipitation was however detected
by electron nicroscopy. It is believed therefore that the age-hardening was entirely due to carbide precipitation, and this was confirmed by the similar age-hardening curves for
both steels, Pigs. 35 a^d 3 6, and also by the same activation
energy for the process, namely 9h-125 KJ/mol (23-30 k cal/mol).
A comparison of the age-hardening intensity against other
carbide forming alloys, Pig. hi* indicates that the increase
in hardness during ageing was less for the high carbon alloy than for the low carbon alloy, as might be expected due to a smaller amount of titanium in solution. The temperature dependence of this maximum increase in hardness during ageing was very different for the two alloys. It can be seen that
the maximum increase in hardness of the 0.li$C steel, which contains much more titanium in solid solution than the
0.35%C steel, did not vary appreciably with temperature until the higher ageing temperatures. This indicates that
at the higher ageing temperatures some TiC precipitation
occurred which was confirmed by electron microscopy, Plates
5 and 6.
The sequence of precipitation in this alloy, was entirely related to the relative diffusion rates of chromium,
molybdenum and titanium and also dependent upon the amount of titanium dissolved in the matrix after solution treatment.
(3 1*6 9,7 5*7 6) early nucleation of ^23^6 Srain boundaries
at lower ageing temperatures was due to the rapid rate of
diffusion of chromium compared with titanium, Plate 5(f)
At higher ageing temperatures the dislocation precipitation
of TiC was preferred due to -a high Ti:C ratio and the smaller
TiC unit cell, which required less surface energy to form a
171