Copy and paste this link to your website, so they can see this document directly without any plugins.


that, with, hardness, materials, after, which, also, higher, content, been, lower, compositions, sections, shown, tool, this, phase, there, high, levels, phases, grain, Powder, Fe-Co-Mo, peak, Metallurgy, however, grade, cutting, hardening


Powder Metallurgy Progress, Vol.13 (2013), No 2 47
H. Danninger, F. Rouzbahani, Ch. Harold, H. Ponemayr, M. Daxelmüller,
F. Simančík, K. Iždinský
Carbon-free precipitation hardened tool steels of the type Fe-Co-Mo/W
have been shown to exhibit as-heat-treated hardness levels similar to
those of carbidic high speed steels but markedly higher resistance to heat
softening. However, the high content of expensive alloy elements is a
marked disadvantage. Here, the effect of the Co and Mo contents,
respectively, on hardness and microstructure has been studied. It has
been shown that the Co:Mo ratio plays a major role, ratios >1 being
preferable; otherwise, the material remains in the α structure during
sintering and heat treatment, respectively, with resulting pronounced
grain coarsening. With lower Co content, the peak hardness is attained at
progressively higher aging temperatures; the absolute hardness level is
however lowered. 15% Mo results in hardness levels that are 5-10 HRC
higher than 10%Mo at the same Co content. Generally it must be stated
that for high speed cutting applications only the materials Fe-25%Co15% Mo and, for slightly less demanding applications, Fe-20%Co-15%
Mo are promising, while lower-alloyed variants are possibly suited for
cold work and hot work applications, in which case their virtually
distortion-free hardening response might be a distinct advantage.
Keywords: Powder metallurgy tool steel, sintering, precipitation
hardening, carbon-free alloys, temper resistance
Tool materials, especially those used for metal cutting applications, are required to
possess high hardness at elevated temperatures to withstand the thermal loading of the
cutting edge [1,2]. In addition to abrasive and adhesive wear, typically resulting in flank
and crater wear, respectively, thermal softening is a typical wear mechanism especially in
high speed tool steels. This is due to the fact that high speed steels are not inherently hard,
as are e.g. hardmetals or cutting ceramics, but can be obtained in a relatively soft, annealed
“manufacturing” state, which makes machining fairly easy, and are then transformed into
the hard “application” state by suitable heat treatment. The hardening mechanisms, in the
case of HSS mainly precipitation hardening through fine secondary carbides, start to fail at
higher temperatures, as a consequence of overaging of the precipitates, leaving only solid
solution strengthening as the remaining mechanism [3]. If the cutting performance of tool
materials is to be improved while the convenient “switch” from soft and machinable to hard
Herbert Danninger, Fardin Rouzbahani, Christian Harold , Technische Universität Wien, Vienna, Austria, and
Materials Center Leoben GmbH, Leoben, Austria
Helmut Ponemayr*, Manfred Daxelmüller, Böhler-Uddeholm Precision Strip GmbH & Co. KG, Böhlerwerk,
Austria, *now: Böhler Profil GmbH, Böhlerwerk, Austria
František Simančík, Karol Iždinský, Institute of Materials and Machine Mechanics, Slovak Academy of Sciences,
Bratislava, Slovak Republic
Powder Metallurgy Progress, Vol.13 (2013), No 2 48
and resistant is to be retained, strengthening mechanisms have to be used that remain
effective up to higher temperatures.
One approach is the use of intermetallic phases as secondary hardening phase in
place of carbides, these phases being less prone to overaging. This has been used for a long
time in hot work tool steels which, in addition to carbides, also contain intermetallic phases.
There is however no need for carbides: As early as 1932, Köster and Tonn [4-6] have
shown that carbon-free alloys Fe-Co-Mo and Fe-Co-W, at that time primarily intended for
hard magnetic applications, can be precipitation hardened to hardness levels >65 HRC. In
the 1960s, Geller and co-workers [7-9] studied numerous carbon-free steel grades; they
found that the absence of austenite stabilizers such as Ni is beneficial for temper resistance.
Fe-Co-Mo and Fe-Co-W-Mo steel grades proved to be particularly well suited for
machining of Ti alloys, being decidedly superior to e.g. standard HSS grade T1. Although
Köster had proposed powder metallurgy manufacturing already in the 1930s [10], it took
many decades until this approach was experimentally followed. In the late 1990s, Karpov et
al. produced Fe-Co-W-Mo tool steel through powder metallurgy, using coprecipitated and
coreduced powders [11]. Danninger et al. [12] showed that tool materials with excellent
properties can be obtained also from elemental powder mixes if the sintering parameters are
adjusted accordingly to result in homogeneous distribution of the alloy elements [13-15]. In
particular the grade Fe-25%Co-15%Mo exhibited an excellent combination of workability
and machining performance. Hardening is afforded by secondary precipitates of the µ phase
(Fe,Co)7(Mo,W)6 that are generated at temperatures >350°C (see also [16-18]). However,
also some µm-sized µ phases should remain during solution treatment to prevent excessive
grain growth and resulting embrittlement.
One pronounced advantage of these tool steels compared to carbidic HSS is their
relatively low hardness as-quenched, which enables machining and even, to some degree,
cold working. Tools can be soft machined after quenching, the final hardening being an
isothermal aging process at moderate temperatures, without any phase transformation.
Thus, geometrical precision after hardening is much better than after the common quenchand temper treatment of carbidic HSS. Recently, the Fe-Co-Mo grade has been made
commercially available by Böhler Uddeholm, under the designation “MC-90 Intermet”
[19], and the applications are mainly tools for which geometrical precision is a must while
hard machining should be avoided as much as possible.
A clear disadvantage of the carbon-free tool steels is however the high content of
expensive alloy elements, in particular with regard to the soaring metal prices in the last
years. In this work it was studied if also lower-alloyed variants of the Fe-Co-Mo grade
might be feasible for producing effective cutting tools with possibly improved workability,
peak hardness and aging response being taken as criteria.
The starting powders used were Carbonyl iron (BASF grade CN), Co powder
(Umex 5-M), and elemental Mo (Plansee, <32 µm). The powders were dry blended for 60
min in a tumbling mixer and then uniaxially compacted at 400 MPa under die wall
lubrication to bars of 55 x 10 x 15 mm and 100 x 12 x 15 mm, respectively. Die wall
lubrication was performed since it had been found that when using admixed pressing
lubricant, the high Mo content results in carbon pickup, with resulting loss of temper
resistance. The compacts were then sintered in a pushtype furnace in flowing hydrogen of
technical purity. Since previous experiments had shown that 2 h sintering at 1370°C had
resulted in homogeneous microstructures with all materials except those containing
>15%W, this sintering regime was chosen also here.
Powder Metallurgy Progress, Vol.13 (2013), No 2 49
After sintering, the density was measured through water displacement, and
metallographic sections were prepared by standard techniques. However, a particular
feature of the materials was the retaining of grinding marks that were concealed after
polishing but reemerged after etching, being visible as parallel lines. Therefore, after the
first polishing step intense overetching was done to remove the deformed surface area, with
subsequent careful diamond repolishing; after this treatment, with most materials clear
sections without grinding marks were obtained after etching (see sections below). In all
cases Picric acid was used for etching. For control purposes, the hardness was measured
also in the as-sintered state.
Then the bars were heated to 1150°C in flowing N2 and hot rolled in 6 passes,
resulting in a total thickness reduction of about 50%; thus the residual porosity could be
eliminated. The specimens were then solution treated at 1150°C for 30 min, oil quenched, and
tempered for 60 min at varying temperatures. Rockwell hardness HRC was measured on cross
sections after carefully cutting the bars using a Labotom saw with water cooling. Selected
specimens were metallographically prepared using the techniques described above.
Thermodynamic calculations
For the ternary system Fe-Co-Mo, there are not too many references in the literature.
In part even fairly recent books and databases show results obtained by Köster and Tonn
many decades ago, apparently since there are no newer data available [20,21]. In Figure1 the
isothermal section at 1300°C is shown, which gives some indication of the phases at sintering
temperature, and the compositions selected for the experiments are indicated. It is evident that
increasing both the Co and Mo levels in parallel does not shift the α-γ transformation too
much; the exact position of the boundary lines is however questionable.
Fig.1. Isothermal section of the system Fe-Co-Mo at 1300°C ([20], after [4]); the
compositions selected for the experiments are plotted.
The system was also calculated using Thermocalc™ software. The first approach,
using the database TCFE, was not successful; the results obtained showed a very poor
correlation with experimental results. E.g. according to these results there should be no
austenite phase at all for Fe-25%Co-15%Mo, which, as shown by the experiments,
definitely exists. Therefore, another database, Kaufman binary, was used, which resulted in
better agreement with experiments, at least qualitatively, although e.g. the solidus
temperatures are given too low. In Figures 2a, b the polythermal sections for 15% Mo and
for 25% Co, respectively, are shown. The austenite-stabilizing effect of Co is clearly
visible, as is the fact that the stability of the µ phase is not too much affected by the α-γ
Powder Metallurgy Progress, Vol.13 (2013), No 2 50
transformation. Co seems to slightly lower the solubility of Mo. From the calculations it
can be deduced that lowering the Co level should stabilize the alpha phase – with possibly
beneficial effect on sintering – and increase the solubility of Mo, resulting in lower volume
fraction of precipitates and thus maybe in slightly lower hardness. Lowering the Mo content
would result in less precipitates and thus in lower peak hardness. Since however the
quantitative agreement between calculation and experiment was still not quite satisfactory,
the exact effects of changing the composition is hardly predictable, and experimental proof
was necessary anyhow.
15%Mo, varying Co content 25%Co, varying Mo content
Fig.2. Polythermal sections of Fe-Co-Mo, calculated through ThermoCalc with database
Kaufman; compositions studied in the present work are marked
As-sintered properties and microstructures
In Table 1 the compositions of the materials tested are given as well as the
resulting density / residual porosity and as-sintered hardness data. Regarding the porosity it
must be considered that this has been calculated using a theoretical density obtained
through the rule of mixture, which is not necessarily applicable here. Nevertheless is it
evident than in all cases near full density has been attained, those materials with a Co:Mo
mass ratio of or below 1 exhibiting slightly lower porosity.
Tab.1. Properties of sintered steels Fe-x%Co-y%Mo, compacted at 400 MPa, sintered 2 h
1370°C in H2
Sintered density
Fe-25Co-15Mo 8.20 + 0.01 2.2 + 0.1 37.4 + 0.7
Fe-25Co-10Mo 7.98 + 0.01 3.7 + 0.1 41.0 + 0.8
Fe-20Co-15Mo 8.20 + 0.01 1.7 + 0.1 39.5 + 0.6
Fe-20Co-10Mo 8.02 + 0.01 2.7 + 0.1 35.6 + 0.1
Fe-15Co-15Mo 8.23 + 0.01 0.7 + 0.1 42.1 + 0.3
Fe-15Co-10Mo 8.02 + 0.05 1.6 + 0.6 36.7 + 0.3
Fe-10Co-15Mo 8.15 + 0.01 1.1 + 0.1 36.8 + 1.4
Fe-10Co-10Mo 7.99 + 0.00 1.7 + 0.0 (5.7 + 0.9)
The hardness levels are in the range of 25 to 42 HRC, there is however no defined
pattern with regard to composition or porosity, i.e. “as-sintered” is similarly undefined as
Powder Metallurgy Progress, Vol.13 (2013), No 2 51
“T1” for sintered aluminium. Since however the materials are not worked or used in this
state, the erratic HRC values are not really of practical relevance.
In Figure3 the as-sintered microstructures are shown. Evidently there are two
distinct groups of morphologies, depending on the composition: if cCo > cMo (Fig.3a-e),
there is the typical microstructure described e.g. in [12, 15] with complex-shaped, heavily
interlocked grains. In the case of cCo = cMo or cCo < cMo (Fig.3f-h) simple polygonal grains
are found, and in part quite pronounced layers of precipitates can be observed at the grain
boundaries, while at the Co-rich variants in part no grain boundary phases are present at all,
in part they are thin and discontinuous, as typically shown in Fig.4. This can be taken as an
indicator that in the ferrite, lattice diffusion of Mo to the grain boundaries to form
intermetallic phases there is markedly faster than in the austenite phase. Quite similar
effects have been observed with Co-free Fe-Mo alloys used for the wear loaded section of
automotive rocker arms, which are subsequently carburized [24]; also in this case,
polygonal grains with continuous precipitates at the grain boundaries have been obtained
after sintering.
25Co-15Mo 25Co-10Mo 20Co-15Mo 20Co-10Mo
15Co-10Mo 15Co-15Mo 10Co-15Mo 10Co-10Mo
Fig.3. Metallographic sections of Fe-x%Co-y%Mo, compacted at 400 MPa, sintered 2 h
1370°C in H2.
Taking the isothermal sections of the system Fe-Co-Mo from [20] it can be
concluded that with all compositions that are located in the γ or α-γ fields, the interlocking
grains are visible while those compositions in the α field exhibit the polygonal
microstructure. This also agrees with the porosity data which for materials sintered in the α
field (see Fig.1) exhibit slightly lower porosity, due to the well known activating effect of
“alpha sintering” [25].
Powder Metallurgy Progress, Vol.13 (2013), No 2 52
25Co-15Mo 15Co-15Mo
Fig.4. Intermetallic phases at grain boundaries in differently alloyed materials.
Heat treatment response
In Figure 5, the isothermal sections of the ternary system are shown for the
temperature “window” suited for solution annealing, the compositions investigated being
plotted once more in the diagrams. It is evident that at 1200°C all the materials are in the
homogeneous α or γ or in the two-phase α-γ phase fields, but there is no µ phase stable any
more. This means that the annealing temperature should be chosen slightly lower, in order
to retain at least some of the µm-size µ phases that prevent excessive grain growth.
Therefore, the intermediate temperature of 1150°C was selected for the experiments.
Section at 1100°C ([20], after [22]) Section at 1200°C ([20], after [23])
Fig.5. Isothermal section of the system Fe-Co-Mo at 1100 and 1200°C; the compositions
selected for the experiments are plotted.
In Figures 6a and b, temper charts are shown for the different compositions after
solution treatment at 1150°C and oil quenching. Figure 6a depicts the materials containing
10% Mo and varying amounts of Co, and Fig.6b shows the same for 15% Mo. It stands out
clearly that for both Mo levels the as-quenched hardness is quite low, at or below 40 HRC,
which enables soft machining. The peak hardness increases with higher Co content and the
onset of precipitation hardening is shifted to lower temperatures while the drop of the
hardness at higher temperatures is hardly shifted, i.e. the range of aging temperatures at
which high hardness levels are attained becomes successively broader. The materials with
Powder Metallurgy Progress, Vol.13 (2013), No 2 53
higher Mo content are generally more tolerant in that respect. This means that at least at
technically relevant hardness levels the shifting of the α-γ transformation to higher
temperatures at low Co:Mo ratio, as shown in Fig.2a, does not play a major role here, in
contrast of the findings of Geller et al. for austenite stabilizers such as Ni [7-9] for which
the transformation temperature is clearly related to the temperature of heat softening.
Fig.6a. Aging response of Fe-x%Co10%Mo, solution treated at 1150°C, oil
quenched, tempered 60 min
Fig.6b. Aging response of Fe-x%Co15%Mo, solution treated at 1150°C, oil
quenched, tempered 60 min
When comparing the 10% Mo and 15% Mo grades (see e.g. Figs.7a and 7b), it is
evident that the higher Mo content generally results in higher peak hardness. In addition to
the reference grade Fe-25Co-15Mo also the grade 20Co-15Mo seems to be attractive.
Grades such as 25Co-10Mo exhibit markedly lower hardness as monolithic materials but
might be interesting as matrix materials for particle reinforced MMCs in which the wear
resistance, in particular against abrasion, is increased by addition of ceramic particles [26].
Fe-25%Co-x%Mo Fe-15%Co-x%Mo
Fig.7. Aging response of Fe-x%X-y%Mo as a function of the Mo content
Microstructures at peak hardness
In Figure 8 the microstructures of peak aged materials with constant Mo level of
15% but decreasing Co content are shown at low magnification to give a general
impression of the microstructures. It stands out clearly that also in this state the polygonal
microstructure is visible for materials with Co:Mo ratio <1 while for 20:15 the fine and
very regular microstructure as typical for 25-15 is revealed. Typically this latter material
contains the µm size µ phases which, as stated above, are necessary for preventing
excessive grain growth during solution annealing, while for the materials with polygonal
Powder Metallurgy Progress, Vol.13 (2013), No 2 54
microstructure such phases are not shown, although the Mo content is the same. This
indicates markedly higher Mo solubility in the latter materials, at least at 1150°C, which
disagrees with the phase diagrams shown in [20] but agrees with the calculated polythermal
section shown in Fig.2a, which predicts lower Mo solubility with higher Co content. At the
lower Mo content, also the Co rich variants show relatively coarse grain structures, which
indicates that the lack of sufficient µm-size µ phases has resulted in the austenite grain
growth that had been observed at the standard 25-15 grade at high solution treatment
temperatures [12, 13].
20-15, 700°C 15-15, 700°C 10-15, 700°C 10-10, 750°C
Fig.8. Metallographic sections of Fe-x%Co-y%Mo, solution treated 30 min 1150°C, oil
quenched, and peak aged. Overview (Low magnification).
25-15, 600°C 25-10, 600°C 20-15, 600°C 20-10, 700°C
15-10, 700°C 15-15, 700°C 10-15, 700°C 10-10, 750°C
Fig.9. Metallographic sections of Fe-x%Co-y%Mo, solution treated 30 min 1150°C, oil
quenched, and peak aged. High magnification.
In Figure 9, the microstructures are shown at higher magnification, each material
being aged at the temperature at which maximum hardness has been recorded. It is evident
Powder Metallurgy Progress, Vol.13 (2013), No 2 55
that once more the coarse polygonal microstructure is observed for Co:Mo <1, with hardly
any structures visible within the grains. This means that the coarse microstructure is caused
by annealing in the ferrite range, at which coarsening is much more rapid than in austenite,
and that also the lack of µm size µ phases results in uninhibited grain growth. For the other
materials the structures are much finer, and µ phases are visible which, as previously shown
[12] are quite effective in preventing grain growth during the solution anneal.
The experiments with Fe-x%Co-y%Mo (x = 10 … 25 mass%; y = 10 ..15 mass%)
have shown that through the blended elemental approach, i.e. mixing of elemental starting
powders, uniaxial compaction and sintering at 1370°C, specimens with almost full density
and homogeneous microstructures can be attained. The materials with Co:Mo mass ratio ≤
1 show slightly less residual porosity, while the microstructures are relatively coarse, with
polygonal grains. The other group of materials shows an interlocking grain structure. This
difference can be attributed to the phases during sintering: while in the first group sintering
is done in the alpha range, the second one sinters in mixed α-γ or plain γ phase.
After precipitation hardening treatment, i.e. solution anneal, quenching and aging,
the materials exhibit in part hardness levels up to 65 HRC and high temper resistance.
When the Co content is decreased, the peak hardness appears at higher aging temperatures
while the hardness itself is lowered. The Mo content strongly affects the peak hardness:
15% Mo results in hardness levels that are 5-10 HRC higher than 10%Mo at the same Co
For high speed cutting applications, primarily the materials Fe-25%Co-15%Mo
and, for slightly less demanding applications, Fe-20%Co-15%Mo seem to be attractive; for
the lower-alloyed variants the alloying cost will probably outweigh the performance;
however, at least some of them might be useful as matrix for particle reinforced metal
matrix composites or for hot work application since for all these grades the virtually
distortion-free hardening by isothermal anneal in the α range can be applied.
This work was carried out within the Materials Center Leoben (MCL) and
financially supported by the FFG through the Kplus initiative. Furthermore, the authors want
to thank Dr.D.Caliskanoglu, Kapfenberg, for carrying out the Thermocalc simulations.
[1] ASM Handbook—Properties and Selection: Irons, Steels, and High-Performance
Alloys. Vol. 1. 10th edition. Materials Park OH : ASM International, 1990
[2] Roberts, G., Krauss, G., Kennedy, R.: Tool Steels. Materials Park OH : ASM, 1998
[3] Fischmeister, HF., Karagöz, S.: Met. Trans., vol. 29A, 1998, p. 205
[4] Köster, W., Tonn, W.: Arch. Eisenhüttenwesen, vol. 5, 1932, no. 8, p. 431
[5] Köster, W., Tonn, W.: Arch. Eisenhüttenwesen, vol. 5, 1932, no. 12, p. 627
[6] Köster, W.: Arch. Eisenhüttenwesen, vol. 6, 1932, no. 1, p. 17
[7] Geller, JuA.: Instrumentalniye Staly. Moscow : Metallurgia Publ., 1983
[8] Brostrem, WA., Geller, JuA.: Metallowedeniye i termicheskaya obrabotka metallov,
1966, no. 11, p. 35
[9] Brostrem, WA., Geller, JuA.: Metallowedeniye i termicheskaya obrabotka metallov,
1970, no. 1, p. 35
[10] Köster, W.: Lecture at the 75th birthday of R.Kieffer (1980)
[11] Karpov, MI. et al. In: Proc. Powder Metall. World Congress 1998, Granada. Vol. 3.
Powder Metallurgy Progress, Vol.13 (2013), No 2 56
Shrewsbury : EPMA, 1998, p. 519
[12] Danninger, H., Rouzbahani, F., Harold, C., Ponemayr, H., Daxelmüller, M., Simančík,
F., Iždinský, K.: Powder Metall. Progress, vol. 5, 2005, no. 2, p. 92
[13] Rouzbahani, F.: PhD thesis. Vienna : TU Wien, 2001
[14] Harold, C.: PhD Thesis. Vienna : TU Wien, 2002
[15] Danninger, H., Harold, C., Gierl, C., Ponemayr, H., Daxelmüller, M., Simančík, F.,
Iždinský, K.: Acta Physica Polonica A, vol. 117, 2010, no. 5, p. 825
[16] Eidenberger, E.: Fundamental investigations of the precipitation behavior of Fe25at%Co-9at%Mo. PhD thesis. Leoben : Montanuniversität, 2010
[17] Stergar, E.: Influence of alloying elements on the precipitation behavior of an Fe-25
m%Co-15 m%Mo base alloy. PhD thesis. Leoben : Montanuniversität, 2010
[18] Eidenberger, E., Stergar, E., Leitner, H., Scheu, C., Staron, P., Clemens, H.: Berg- und
Huettenmaenn. Monatshefte, vol. 153, 2008, p. 247
[19] Böhler Uddeholm Data Sheet “MC-90 Intermet”, Kapfenberg (2013)
[20] Raynor, GV., Rivlin, VG.: Phase Equilibria in Iron Ternary Alloys. London : The Inst.
of Metals, 1988
[21] Villars, P., Prince, A., Okamoto, H.: Handbook of Ternary Phase Diagrams. Vol. 6.
Materials Park OH : ASM, 1995, p. 8244
[22] Loo, FJJ. van, Bastin, GF., Vrolijk, JWGA., Hendrick, JJM.: J.Less Common Met., vol.
72, 1980, no. 2, p. 225
[23] Das, DK., Rideout, SP., Beck, PA.: J.Met., vol. 4, 1952, no. 10, p. 1071
[24] Seyrkammer, J., Blaimschein, F., Delarche, C., Pourprix, Y.: Adv. Powder Metall. and
Particulate Materials – 1992 No. 5. Eds. J.M. Capus, R.M. German. Princeton NJ:
MPIF, 1992, p. 141
[25] Lenel, FV.: Powder Metallurgy – Principles and Applications. Princeton NJ : MPIF,
1980, p. 413
[26] Danninger, H., Rouzbahani, F., Harold, C., Ponemayr, H., Daxelmüller, M., Simančík,
F., Iždinský, K.: Int. J. Powder Metall., vol. 45, 2009, no. 6, p. 27

PDF Document reader online

This website is focused on providing document in readable format, online without need to install any type of software on your computer. If you are using thin client, or are not allowed to install document reader of particular type, this application may come in hand for you. Simply upload your document, and Docureader.top will transform it into readable format in a few seconds. Why choose Docureader.top?

  1. Unlimited sharing - you can upload document of any size. If we are able to convert it into readable format, you have it here - saved for later or immediate reading
  2. Cross-platform - no compromised when reading your document. We support most of modern browers without the need of installing any of external plugins. If your device can oper a browser - then you can read any document on it
  3. Simple uploading - no need to register. Just enter your email, title of document and select the file, we do the rest. Once the document is ready for you, you will receive automatic email from us.

Previous 10

Next 10