Hyper baric carburising process - Gruppo Italiano Frattura

PROCESSI
Hyper baric carburising process
E. Gianotti
A new patented technology improves low pressure carburising in vacuum furnace and traditional
carburising in sealed quench furnace.
The author, starting from argument that is possible to see in the specialized publications on the subject,
has examined by scientific criteria which are the limits and the advantages of both the low pressure
carburising ,or the controlled atmosphere process. He evaluates the technological possibility to overcome
some of the limits that are conditioning the mechanical property of the heat treated pieces. Before
suggesting new technological solutions that can promote better metallurgical properties together with
clean environment, he examine some scientific arguments to make clear the problems.
At the end of this research some important technological modification will be realised and the obtained
results confirm the validity of the new process of hyper baric carburisation.
The cause of the defects arising from the traditional technology have been eliminated and the
modifications requested on the furnaces are not expensive. In same cases, maybe the new furnaces are
even cheaper, both in the buying phase than in the running cost.
Memorie
Key words: steel, carburising, heat treatment, plants and fixture
INTRODUCTION
The low pressure carburising has already reached a good
diffusion in the heat treatment market .
Compared with the old traditional controlled carburising atmosphere,
it has the big advantages of being near pollution
free.
Some other advantages claimed by vacuum furnace builders
have some difficulty being demonstrated by practical application,
as it’s possible to see in the specialized publications
on the subject.
Avoiding taking into consideration some less technical arguments
like cost effectiveness, speed of carbon penetration,
and furnace and facility cost , the most stimulating discussions
are about the technological property of the product,
compared with the others obtained with the old controlled
atmosphere.
The most discussed arguments are the fatigue strength and
the toughness that are conditioned by:
- Surface decreasing of alloying elements.
- Inter granular oxidation.
- Speed of cooling in the quench step, using gas or gas mixture
up to 20 bars.
You can read about this subject in the bibliography at the
end of this work 1) , 2) , 3).
Starting from this argument, the author has examined, by
scientific criteria, which are the limits and the advantages of
both the low pressure carburising ,or the controlled atmosphere
process. He evaluates the technological possibility to
overcome some of the limits that are conditioning the mechanical
property of the heat treated pieces.
Before suggesting new technological solutions that can promote
better metallurgical properties together with clean environment,
we need to examine some scientific arguments to
make clear the problems.
a) Solid-gas phases equilibrium of the metals to high temperature.
b) Thermodynamic equilibrium of the surface inter granular
oxidation that can happen during the carburisation with
Elio Gianotti
Trattamenti Termici Ferioli & Gianotti SpA - Rivoli Torino Italy
endogas or in the vacuum furnace if there are too many
vacuum leaks.
c) The quenching in gas, when the heat treatment is made in
vacuum furnace, is too slow, also if it’s made with high
pressure or blending of same expensive gas like He. In
this case the toughness is lower than the quenching in oil.
(See the bibliography 3) by Fernando Da Costan just
quoted).
SOLID-GAS PHASES EQUILIBRIUM OF THE METALS
IN FURNACE TO HIGH TEMPERATURE
Metals evaporation in vacuum furnace
Every metal has a tension, or pressure of evaporation, that is
only temperature dependent (see tab 1.)
In an imaginary vacuum furnace, without leakage, if the absolute
pressure that can reach the vacuum pump is lower
than the evaporation tension of the metal, the metal evaporates
continuously and goes to solidify on the cold walls of the
furnace or of the pump that have a temperature lower than
the metal evaporation. The remaining part of the vapour is
eliminated by the pump.
The metal that solidifies on cold walls decreases the partial
pressure in the furnace, so other metal can evaporate.
In the event where the absolute pressure generated by the
pump is not lower than the metal evaporation tension, the
metal vapour , also if it has saturated the furnace chamber ,
cannot be extracted by the vacuum pump. It can also, in this
event, solidify on the cold walls if these are at the temperature
lower than the solid-gas equilibrium and allows further
evaporation of the metal surface.
In the more realistic event in which the furnace has some
leakage, and if the vacuum pump is not able to reach value
lower than metal evaporation pressure, some metal vapour
can be sucked together with the air from the leakage,
breaking the equilibrium of the ‘solid metal – vapour ’ and
permitting again the evaporation of the metal surface.
Therefore, in every event, the heating of the steel in the vacuum
furnace generates a continuous loss of alloying elements
from the surface. Only in the purely theoretical event
of a vacuum furnace without leakage, with the inner walls at
the same working temperature of the furnace and the vacuum
pump isolated, would it be possible to avoid continuous
evaporation of the alloying elements.
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7-8/2006 Memorie
Metals evaporation in controlled atmosphere furnace
Also in this furnace, the metal’s vapour tension generates an
impoverishment on the steel’s surface that is more important
as the temperature goes higher. The impoverishment will be
higher in the alloying metals that have a higher tension (see
table 1), so the surface composition can vary in its alloying
elements.
The impoverishment in this event is linked to the passage of
carburising gas over the pieces’ surface.
It’s not easy to say which of the two furnaces is more reliable
regarding this problem, because many variables are to
be considered.
The vacuum furnace, with a good seal, surely has a lower
number of atmosphere changes compared with the atmosphere
competitor, but it has a great surface of cold walls
that makes easy the vapour deposition of the alloying metals
and can generate other evaporation.
To avoid evaporation loss of alloying metals, considering
that it is not possible to change the evaporation equilibrium,
it’s only possible to decrease the atmosphere change in the
furnaces and to maintain the walls’ temperatures near the
process temperature.
Utilizing low pressure carburisation by a vacuum furnace,
the process temperature can arrive up to 1.000° C . At this
temperature the Mn is the metal that has the highest vapour
tension, followed in decreasing order by Al, Cu, Cr. It may
be considered that the Mn vapour tension at 1.000°C is
about 10 -2 mbar, while the absolute pressure at the furnace
during the carburising process is about 10-30 mbar, so the
vacuum is not too dangerous.
Table 1 – Vapour tension of
metals in the range 0° to
2800°C.
Tabella 1 – Tensione di
vapore dei metalli a
temperature da 0 a 2800°C.
To better value the dangerous effect of the evaporation of alloying
elements in vacuum furnace and particularly in the
low pressure carburising process, see in the bibliography the
work of B.Clausen 1) and Y. Bienvenu, K Vieillevigne 2).
THERMODYNAMIC EQUILIBRIUM
OF INTER GRANULAR OXIDATION
Carburising process in controlled atmosphere
The most utilized atmospheres are the endogas and the nitrogen-methanol
. In both cases the blend of the gas is formed
by about 20% of CO that is the most important in the
exchange of the C from the atmosphere to the surface of the
carburising steel.
The schematic reaction is the following:
2CO ↔ 2C + O 2
.
The free energy ∆ G of the reaction at 927° C (1200 K) calculated
by the Barin Knacke parameter is
2CO ↔ 2C + O 2
with ∆ G = + 104149 Kcal at 927°.
The constant K of the reaction will be:
lnK = - 104.149 / (1,987 · 1200) = - 43,68 so K = e - 43,68
but K = p (O 2
) / p (CO 2
) 2 than substituting the values will be:
p (O 2
) = e - 43,68 · 0, 2 2 = 4,3 -21 = 10 -20,37
This value of p (O 2
) generates about 1.140 mVolt in the
oxygen probe (see Nerst formula) and is equivalent to a car-
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PROCESSI
bon potential of about 0,80% (see oxygen probes tables).
It is so demonstrated that in the carburising atmosphere there
is a pO 2
equal to near 10 -20 and thereafter every element
that is in such atmosphere, with a chemical equilibrium with
the oxygen partial pressure equal or lower , will become an
oxide.
In table 2 there are some alloying elements that are always
in the steel, also like impurity, with the relative partial oxygen
pressure in equilibrium at 900°C with the element in the
oxidising chemical reaction.
Note as the metals Cr, Mn, Si, Al, normally present in the inter
granular oxidation, oxidise with pO 2
from 10 -24 to 10 -36 .
To avoid their oxidation, it is necessary to operate with pO 2
below 10 -36 .
For more detailed information about the inter granular oxidation
in the atmosphere controlled furnace, see ref.4).
Low pressure carburising process in vacuum furnace
Both in the low pressure carburising process or in plasma
carburising process ( ionic discharge), the working pressure
in the furnaces are about 10 – 200 mbar.
Also, when the process is starting and there is the need to
clean the chamber and the pieces, the higher vacuum reached
by the pump is 10 -4 – 10 -5 bar; that is a vacuum very far
from the 10 -36 /0.21 bar necessary to avoid the inter crystalline
oxidation.
The purity of the furnace is not sufficient, but the continuous
introduction of high purity carburising gas, as N 2
, Ar 2
or hydrocarbon
like C 3
H 8
, CH 4
, C 2
H 2
or other, washes the atmosphere
until it arrives at the needed purity.
Normally, with furnaces with good sealing , it is possible to
arrive at the end of the process, avoiding inter granular oxidation.
The furnace sealing may deteriorate if the maintenance is
not good. In this event, if the furnace absorbs air from the
leakages, the oxygen concentration may increase over 10 -36 ,
and the inter crystalline oxidation occurs.
STIRRING OF CARBURISING ATMOSPHERE
While in the controlled atmosphere furnace it is possible to
homogenize easily the atmosphere composition in the inner
furnace by one or more fans, it is more difficult to avoid
stratifications or non uniformity in the vacuum furnace.
There are some systems realized, the most curious has been
made up by an important furnace builder in U.S. He has
mounted a big fan in the furnace chamber.
The non uniformity in the atmosphere may generate hardness
and case depth problems.
Réactions P(O 2
) [bar] ∆G [Kcal]
2Fe + O 2
⇔ 2FeO 10 -16,7 -88
4/3 Cr + O 2
⇔ 2/3Cr 2
O 3
10 -24 -130
2Mn + O 2
⇔ 2MnO 10 -27 -145
Si + O 2
⇔ SiO 2
10 -30 -160
4/3 Al + O 2
⇔ 2/3 Al 2
O 3
10 -36 -200
2Mg + O 2
⇔ 2MgO 10 -44 -230
2H 2
+ O 2
⇔ 2H 2
O 10 -17 -90
2CO + O 2
⇔ 2CO 2
10 -16 -87
Table 2 – Partial pressure p(O 2
) in equilibrium at 900°C with the
metals oxide and the CO and H 2
.
Data obtained from Ellingham-Richardson diagram.
Tabella 2 – Equilibrio a 900°C della p(O 2
) con gli ossidi dei
metalli elencati e con l’idrogeno e il CO.
Valori ricavati dal diagramma di Ellingham-Richardson.
THE NEW HYPER BARIC TECHNOLOGY ELABORATION
After we have seen all the advantages and the lacks of the
old technologies, it is possible to engineer a new technology
that takes into account only the advantages and eliminates
the defects.
In the low pressure carburisation, the problems are:
a) Evaporations of alloying elements determined by low
pressure and cold walls
b) Few homogeneity of carburisation in the different zone of
the furnace.
c) Not satisfactory toughness due to slow cooling in the gas
quenching .
In the controlled atmosphere carburisation the problems are:
d) Environmental, because of the pollution of atmosphere
gas and quenching oil
e) Inter crystalline oxidation with case depth proportional to
square root of time, due to O2 existing in the endogas.
f) Evaporations of alloying elements determined by carburising
atmosphere exchange.
Vacuum furnace modifications
It’s possible to avoid the a) and b) defects by modifying the
furnace so that it can work above the atmospheric pressure.
The cold wall will be thermally insulated using ceramic fibre
or other light insulating materials and inner this insulation
will be placed a strong, atmosphere tight crucible that
also has the task of sustaining the carburising load.
It will be necessary to change the instruments for the vacuum
regulation by instruments suitable for regulation of
overpressure.
The vacuum pump may be eliminated otherwise, if existing
may be utilised to clean more rapidly the furnace atmosphere.
After five or more atmosphere changes by cryogenic nitrogen,
with purity of 5 ppm O 2
, the furnace will be held in light
pressure , normally 500 mm of water. This way, it will
avoid air entrance by possible leakage.
The carburising atmosphere must be agitated by a fan to
guarantee the carburisation homogeneity.
To avoid the problem c) about the insufficient toughness, the
unique solution is the oil quenching, so it is necessary to
make provisions for an oil tank below the atmosphere quenching
chamber. This is particularly useful for the commercial
heat treatments that sometimes have many chemical
steel compositions in the same load.
Atmosphere furnace modifications
The furnace must be modified with an addition in the inner
part of a strong, atmosphere tight, crucible that has also the
task of sustaining the carburising load, both in heating
chamber or in the quenching zone. Also, the over pressure
security valve must be gas tight.
After this modification, the furnace can work as described in a
vacuum modified furnace, with a tight nitrogen atmosphere.
This is the condition to avoid the problem scheduled in d) and
f), determined by the flow of the endothermic atmosphere.
ELIMINATION OF INTER CRYSTALLINE OXIDATION
Examining tab. 2, it is possible to note that for avoiding the
alloying element oxidation, that is the origin of the inter granular
oxidation, it is necessary to keep the oxygen partial
pressure below 10 -36 .
With the modification made with respect to the traditional
furnaces , both vacuum or traditional controlled atmosphere
, it is possible to arrive at a sealed chamber, delimitated by a
crucible made in refractory steel, with leakage almost inexistent
and at worst with the gas that goes out and not with the
air that enters to damage the carburising atmosphere.
Memorie
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7-8/2006 Memorie
Fig. 1 – Hyper baric sealed
quench furnace featured with
the most important
modifications compared to
vacuum or controlled
atmosphere furnace.
1) Water cooled sealing
gasket in the outer door.
2) Over pressure bleeder with
sealing, temperature resistant.
3) Stirring fan for carburising
atmosphere.
4) Quenching chamber with
oil or pressurized gas.
5) Computerized control of
nitrogen during the cleaning ,
the carburising, and the
quenching time.
6) Inner door.
7) Carburising computerized
control, according to the
boost-diffusion cycle, with
hydrocarbon like propane,
cyclohexane alcohol.
Fig.1 – Impianto a camera
per la cementazione
iperbarica con evidenziate le
principali modifiche rispetto
ai forni a vuoto e ad
atmosfera controllata.
1) Guarnizione di tenuta atmosfera , raffreddata ad acqua , per la porta anteriore
2) Sfioratore di sovra pressione, con guarnizione a tenuta termica, per permettere l’uscita di gas caldi o fiamme.
3) Ventilatore di omogeneizzazione nella camera di carburazione coibentata.
4) Sistema di raffreddamento in azoto nella precamera fredda.
5) Sistema computerizzato di controllo e regolazione della pressione dell’azoto durante il lavaggio, la carburazione e l’eventuale tempra in
azoto.
6) Porta intermedia del forno.
7) Sistema computerizzato per la cementazione secondo il ciclo boost-diffusion con idrocarburi ( propano, acetilene, alcool cicloesano ).
The atmosphere composition in the carburising crucible is
cryogenic nitrogen with about 5 ppm O 2
that is O 2
= 5.10 -6 .
The purity of atmosphere is not sufficient to avoid inter granular
oxidation, also in the event of many changes of nitrogen
before starting the carburising.
In the furnaces derived from vacuum technology, the cleaning
may be accelerated using the vacuum pumps, and without
them may be utilized the formula referred to point 5 of bibliography).
The furnace is normally atmosphere tight, and possible
leakage doesn’t permit air to enter, but only nitrogen to go
out from the crucible. In this case some other nitrogen will
be introduced to restore the hyper baric pressure.
If the furnace atmosphere purity is the same as the cryogenic
nitrogen, that is 5.10 -6 , to arrive at the needed value of 10 -36 ,
it is necessary to eliminate still O 2
. For example, an industrial
furnace with a crucible volume of about 1 m 3 , equal to
44,64 gas moles, contains about 44.64·32·5·10 -6 = 0.0071 g
of O 2
. In the event of oxygen concentration equal to 10 -36
the O 2
content will be 44.64·32·10 -36 = 1.428·10 -33 g of O 2
.
So it is necessary to eliminate 7 mg of O 2
from the furnace
atmosphere to be sure to avoid the inter granular oxidation.
Only for curiosity, utilising the Avogadro number N A
=
6.02.10 23 mole –1 , it’s possible to calculate that a concentration
of O 2
= 10 -36 corresponds to about 6.02.10 -13 molecule
of O 2
every 22.4 litre of gas, that is one molecule of O 2
every (22.4 / 6.02) 10 13 litre, or 3.7.10 10 m 3 of gas. This value
is so little that is impossible to measure it.
To overtake this problem, there is an innovative idea of the
hyper baric carburisation. To put in the atmosphere tight
crucible with the load, a thin metal shaving or metal sponge
with an equilibrium pressure of oxidation to 1,000°C below
10 -36 volumes of O 2
.
One metal that is easy to find with this characteristic is the
Fig. 2 – Hyper baric
pit furnace featured
with the most
important
modifications
compared to
controlled
atmosphere pit
furnace.
In this case the
quenching is made
out of the furnace.
Fig. 2 – Impianto a
pozzo per la
cementazione
iperbarica con
evidenziate le più
importanti modifiche
rispetto al forno a pozzo ad atmosfera controllata.
In questo caso la tempra dovrà essere fatta in una vasca esterna
al forno.
Titanium (see tab 2).
The shaving or the sponge must be very light, and must be
loaded in the furnace with the pieces to be carburised. The
free energy (∆G) of the chemical reaction between Ti and O 2
is lower than that of the alloying element of the steel, so that
the most part of atmosphere oxygen is captured by the Ti,
and avoids forming inter granular oxidation.
The weight, or better the surface of the Ti sponge or shaving
must be sufficient to react with all the oxygen in the atmosphere,
therefore it is important, more than the weight, that
can be ten - twenty grams, the surface extension of the Titanium
.It is best to put it near the fan where there is a good
agitation of the atmosphere.
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THE HYPER BARIC CARBURISING PROCESS
The carburisation can start when the crucible is full of clean
cryogenic nitrogen and in pressure by introducing hydrocarbons,
like that utilized in the low pressure process.
The quantity of hydrocarbons, the time of introduction, and
the time of diffusion are the same as the boost-diffusion process
in vacuum furnace. The technology of carbon enrichment
is the same because it is not possible to analyze the atmosphere
like in the controlled atmosphere furnace.
It is possible to theoretically foresee the boost and diffusion
time utilizing some algorithm obtained from practical experiment
(see ‘Algorithm for carbon diffusion computation in a
vacuum furnace’ ref. 6)) . Moreover, it is easier than in the
vacuum furnace to take out the furnace samples to control the
surface carbon content by spectrometer and the case depth.
It is correct to remember that the oxygen probe is not suitable
to control the inter crystalline oxidation because it’s not
sensible enough. The highest signal utilizable with confidence
is 1,200 mV. Applying the Nerst formula:
mV = 0.0496·K·[logO 2
(air) - log O 2
(furnace)].
It is possible to calculate the minimum partial pressure of O 2
that can read the oxygen probe for a temperature of 1.200K
that is 927°C:
mV 1200 = 0.0496·1200 (-0.678-log x )
calculating x = -20,84,
therefore pO 2
= 10 -20.84 .
The oxygen probe may be therefore only utilized like an
alarm sensor, to signal some anomalous situation in the atmosphere
or in the cleaning of the crucible.
In the fig. 3 there is a picture of the little laboratory furnace
utilized for the first carburising tests with cryogenic nitrogen
with 5 part for million of O 2
, Titanium shaving and propane.
The crucible volume is about 1,000 cm 3 .
Before starting the steel sample to be carburised is loaded in
the crucible that is successively sealed. Then the crucible is
washed by five change of nitrogen atmosphere. After the last
change, the pressure in the crucible was put 500 mm of water
column.
At this point can start the heating cycle. When the furnace arrive
in temperature the propane can be introduced.
The boost period was 30 minutes and the diffusion time 2
hours. No addition of nitrogen has been necessary during the
2.30 h of the whole cycle.
The temperature was 950°C and the sample 18 Ni Cr Mo 4
steel. The final results have been: surface carbon content
0.80%; case depth 0.55 mm; inter crystalline oxidation absent.
CONCLUSIONS
The obtained results confirm the validity of the new process
of hyper baric carburisation. The cause of the defects arising
from the traditional technology have been eliminated and the
modifications requested on the furnaces are not expensive.
In same cases, maybe the new furnaces are even cheaper,
both in the buying phase than in the running cost.
The most innovation in the new technology may be so summarized:
1) Compared to vacuum furnace
No necessity of the vacuum or low pressure with consequent
elimination of the vacuum pump and the leakage problems.
Possibility to set up a fan in the carburising chamber for better
uniformity in surface carbon and case depth.
Introduction of Titanium shaving or sponge with the carburising
load, The high affinity of Titanium with the oxygen
avoids totally the inter crystalline oxidation, while in the
low pressure process, it is not always so.
2) Compared to controlled atmosphere furnace
The hyper baric pressure in the furnace allows avoiding contamination
of the carburising atmosphere by the air through
the leakage.
The saving of a continuous flow of endogas make the furnace
more similar to the vacuum, whether for flexibility or for
environment protection.
The introduction, with the carburising load, of Titanium shaving
or sponge, avoids totally the inter crystalline oxidation,
while in the controlled atmosphere that is not possible.
BIBLIOGRAPHY
1) Détérioration de la couche superficielle après cémentation
basse pression. B. Clausen, F.Hoffmann, P. Mayr
(Traitement Thermique n° 355 Mai 2004).
2) Influence des réactions gas-solides lors du traitement
thermique des métaux a haute température et sous pression
réduite. Yves Bienvenu et Karine Vieillevigne
(Traitement Thermique n° 362 Avril 2005).
Memorie
3) Comparaison de différents types de refroidissement
après traitement thermochimique et influence sur les caractéristiques
mécanique et métallurgiques. Fernand Da
Costan (Traitement Thermique n° 357 Aout-Sept.
2004).
4) Oxidation intercristalline lors de la cémentation gazeuse
avec endogas. (Grain boundary oxidation in endothermic
gas carburising process). Elio Gianotti, Irene Calliari,
Marzia Zanesco, Emilio Ramous, DIMEG Università di
Padova/Italia.
Fig. 3 – Laboratory furnace with device for hyper baric
carburisation.
Fig. 3 – Fornetto di laboratorio attrezzato per le prove di
carburazione iperbarica.
5) Calcolo del volume di gas necessario per il lavaggio dei
forni ad atmosfera controllata. Elio Gianotti (La Metallurgia
Italiana n°4 anno 2004).
6) Algorithm for carbon diffusion computation in a vacuum
furnace. Experimental methods predict carburising times.
Elio Gianotti (Heat Treating Progress. November
2002 ).
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7-8/2006 Memorie
CEMENTAZIONE CARBURANTE IPERBARICA
Parole chiave:
acciaio, carburazione,
trattamenti termici, impianti e attrezzature
Una nuova tecnologia (patented) che migliora la cementazione
in bassa pressione nei forni a vuoto e la cementazione
in gas nei sealed quench.
La diffusione a livello industriale della cementazione a bassa
pressione si è ormai ritagliata una fetta di mercato. Essa
presenta l’indubbio vantaggio sull’atmosfera controllata
d’essere meno inquinante. Altri vantaggi vantati dai costruttori
sono di meno facile riscontro come dimostra molta pubblicistica
nata sull’argomento.
Tralasciando alcuni argomenti di carattere più commerciale
come costi di produzione, velocità di carburazione, costi degli
impianti, le discussioni più stimolanti riguardano le caratteristiche
tecnologiche del prodotto rispetto a quelle ottenute
con la vecchia tecnologia dell’atmosfera controllata.
Gli argomenti più controversi riguardano la resistenza a fatica
e la resilienza, condizionate da:
- Impoverimento superficiale di elementi leganti.
- Ossidazione intercristallina.
- Drasticità di tempra con miscele varie di gas e pressioni
ABSTRACT
fino a 20 bar.
Si veda a questo proposito la bibliografia al fondo di questa
memoria ai numeri 1), 2), 3).
Traendo spunto da questi argomenti si è voluto esaminare
da un punto di vista scientifico quali sono i limiti ed i vantaggi
sia del processo in bassa pressione che di quello iperbarico
e valutare le possibilità tecnologiche del superamento
di alcuni di questi limiti che condizionano le caratteristiche
meccaniche del prodotto cementato e temprato.
Per proporre delle soluzioni tecnologiche nuove che uniscano
ai vantaggi di un processo ecologicamente “pulito” delle
caratteristiche metallurgiche migliori vengono quindi dapprima
esaminati:
A Gli equilibri delle fasi solido – gas dei metalli ad alta
temperatura.
B Gli equilibri termodinamici dell’ossidazione intercristallina
superficiale che avviene durante la cementazione
con endogas o nei forni a vuoto che hanno perdite di
vuoto eccessive.
C I problemi che nascono dalla tempra in gas che essendo
meno drastica di quella in olio, malgrado tutti gli accorgimenti
fino ad oggi inventati, genera un decadimento
delle caratteristiche meccaniche compresa quella particolarmente
critica della resilienza. Vedasi studio di Fernand
Da Costan 3) già citato.
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TRATTAMENTI TERMICI
A new approach to hardening mechanisms
in the diffusion layer
of gas nitrided α-alloyed steels.
Effects of chromium and aluminium:
experimental and simulation studies
C. Ginter, L. Torchane, J. Dulcy, M. Gantois, A. Malchère, C. Esnouf, T. Turpin
Memorie
Hardening mechanisms in the diffusion layer of gas nitrided α-iron and -steels have been investigated
through the study about effects of chromium (binary alloys and industrial steels) and aluminium
(industrial steel). After nitriding (520°C 48h), nitrogen mass balance between total nitrogen
concentration located in the diffusion zone, experimentally determined, and the expected theoretical
nitrogen concentration, reveals for each alloy a “nitrogen excess”. Jack and Mittemeijer [1-3] suggested
that the volume misfit between semi-coherent nitrides and matrix induces local matrix lattice distorsion,
leading to a local increase of nitrogen solubility in the matrix.
We propose a new approach, based on thermodynamical calculations (Thermo-Calc software), confirmed
by different characterization methods (HRTEM, EDX and X-Ray). Indeed no significant solid solution “N
excess” occurs, but the total nitrogen concentration is explained by complex MN nitrides precipitation,
isomorph of CrN FCC, containing chromium, iron (up to 30at.% at 50µm from the surface), molybdenum
and vanadium. During annealing (520°C 48h), atomic iron fraction in MN nitrides decreases and the
corresponding nitrogen atomic fraction diffuses to the core.
Addition of aluminium in industrial steel strongly increases nitrogen concentration and hardening
(∆=HV x
-HV initial
). Aluminium induces in the diffusion layer precipitation of Fe 4
N and Fe 2-3
N and
precipitates in complex MN FCC nitrides, containing chromium, iron and molybdenum.
Key words: Nitriding, Nitrogen Excess, Nitrides, Chromium, Aluminium, Hardening mechanisms
INTRODUCTION
Gas-nitriding is a thermochemical treatment, applied in aeronautic
and automotive industries, to improve the fatigue
resistance, tribological and anticorrosion properties. The nitrided
case can be divided, as a general rule, into a compound
layer adjacent to the surface (thickness about 10µm)
and a diffusion zone (depth up to 1µm). This study concerns
only the diffusion zone. Relationship between nitrogen concentration
and hardness are strongly dependant of steel composition
and microstructure: different hardening mechanisms
occur, depending of nitriding time and temperature.
The purpose of this paper is to propose a new explanation
about nitrogen localization and improve hardening mechanisms
understanding. Thus effects of chromium and aluminium
were investigated through experimental, simulation
and characterization studies.
Caroline Ginter
Laboratoire de Sciences et Génie des Surfaces (LSGS), Ecole des Mines, Nancy, France
Aubert et Duval, Eramet Group
L. Torchane, J. Dulcy, M. Gantois
Laboratoire de Sciences et Génie des Surfaces (LSGS), Ecole des Mines, Nancy, France
Annie Malchère, C. Esnouf
Laboratory GEMPPM, CECM Group, INSA Lyon, France
T. Turpin
Aubert et Duval, Eramet Group
Paper presented at the 2 nd International Conference
HEAT TREATMENT AND SURFACE ENGINEERING IN AUTOMOTIVE APPLICATIONS
organised by AIM, Riva del Garda, 20-22 June 2005
Binary and ternary alloys have been the subject of many investigations
[1,2,4,5,6], which showed that aluminium and
chromium increase nitrogen content and hardness, due to
the chromium nitrides FCC semi-coherent precipitation. Nevertheless
mechanisms due to aluminium are not clear [7-9].
Moreover hardening mechanisms in industrial steels are
complex and at this time not really understood, in spite of
two important studies about 32CrMoV13 steel [10,11], which
suggested interesting mechanisms related to the chromium
presence.
Therefore it was first necessary to determine nitrogen localization
on binary alloys, and explain the “N excess” observed
by numerous authors in binary and ternary alloys
[1,2,12,13,14]. Then, using the same methodology (experiment,
simulation, characterization) on industrial alloys, nitrogen
localization identification in industrials steels allowed
to suggest several hardening mechanisms, related to
chromium and aluminium presence.
MATERIALS AND EXPERIMENTAL PROCEDURE
Materials investigated in this work are binary alloys (1, 3
and 5wt.%Cr) and about 14 industrial α-alloyed steels manufactured
by Aubert et Duval (heated, quenched then annealed).
Two industrial steels were specially studied; their
chemical composition is given in Table 1.
Samples were gas-nitrided in a vertical furnace during 48h
at 520°C, then quenched in water. The thermogravimetric
measurements were useful to follow the nitrogen mass tran-
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TRATTAMENTI TERMICI
7-8/2006 Memorie
Steel wt.%C wt.%Si wt.%Mn wt.%Ni wt.%Cr wt.%Mo wt.%V wt.%Al
Binary Alloys / T6 wt.%Nexp wt.%N Excess
wt.%N LCth
520°C Kn1 48h
50µm surface 50µm surface
Fe - 1wt.%Cr 0.27 0.44 0.13
Fe - 3wt.%Cr 0.80 1.13 0.29
Fe - 5wt.%Cr 1.35 1.84 0.45
HV50
Initial
K 0.318 0.287 0.532 0.09 3 0.83 0.27 0.013 329
I 0.311 0.330 0.560 0.13 1.73 0.30 0.94 296
sfer through the material from the gas/solid interface. Some
of them were then annealed, in N 2
atmosphere, during 48h at
520°C, then quenched in water.
Nitrided and annealed samples were mechanically characterized
with a micro-hardness tester. Nitrogen profiles in the
nitrided zone were determined from the sample surface through
an electron microprobe analysis. Phases identification
was carried out using X-Ray θ-2θ diffractometer, employing
CoKα radiation.
Nitrides and carbides were examined through Transmission
Electron Microscopy, using a Jeol 200CX (200kV) for conventional
imaging and a Jeol 2010F with a field emission
gun for high resolution imaging and nanoanalysis, equipped
with an INCA Energy Dispersive X-Ray System from
Oxford, with polymer ultra-thin window. TEM observations
were performed on extraction replicas and thin foils. Extraction
carbon replicas were collected on copper grid, after carbon
film deposition on a polished surface and nital
(1.4%HNO 3
+ ethanol) chemical etching. Thin foils were
prepared by spark eroding discs of 3mm diameter (initially
about 700µm in depth), then mechanically polished with a
tripod polisher up reaching 20 to 50µm in depth, to be finally
ion thinned.
Thermodynamical calculations were performed with Thermo-Calc
software, based on Gibb’s free energy minimization
of the defined system. Phases (molar fraction with atomic
chemical composition) appearing at the chemical equilibrium
were simulated, from experimental data (nitriding
temperature, alloy chemical composition, nitrogen concentration).
Table 1 – Chemical
composition and initial
hardness of K and I industrial
steels.
Tabella 1 – Composizione
chimica e durezza iniziale
degli acciai industriali K e I.
I - Binary alloys - Chromium effect
It is well known that nitrogen concentration and hardness
profiles after nitriding depend on chromium concentration.
Indeed, due to the strong Cr-N affinity, chromium combines
with nitrogen to precipitate as fine semi-coherent nitride
CrN FCC, directly responsible for observed hardening.
As hardening mechanisms in industrial steels revealed really
complex, a preliminary work was conducted on binary alloys.
It should be noticed that a discontinuous precipitation,
already observed by several authors [3,11,15,16,17,18,19],
occurs in Fe-3wt.%Cr (Fig.1) and 5wt.%Cr alloys, leading to
a hardness decrease (about 200HV) and does not occur in
steels. Consequently, the objective of this work on binary alloys
was concentrated on nitrogen localization and role. The
Fe 4
N phase often grows at grain boundaries at the compound
and diffusion layers interface, up to 10 or 20µm at the beginning
of the diffusion layer. Consequently, it was decided to
investigate the diffusion layer from 50µm below the surface.
Three binary alloys (1, 3 and 5wt.%Cr) were nitrided at
520°C during 48h with Kn = P(NH 3
)/P(H 2
) 3/2 of 1, then
mass nitrogen concentration profiles were performed (Fig.
1). Nitrogen mass balance was determined for each alloy as
follows: N theoretical, called from now LCth (mass nitrogen
Limited Concentration), subtracted from total amount of N
absorbed [N total
], experimentally determined, revealed N excess
for each alloy below the surface. LCth is defined as the
nitrogen mass concentration involved in nitride precipitation
CrN [N CrN
] and the nitrogen mass concentration dissolved in
the matrix [N α
] (0.043wt.%N at 520°C [20]). LCth depends
on nitriding temperature and alloy chemical composition
Fig. 1 – Mass Nitrogen
concentration profiles versus
nitriding depth of Fe - 1, 3 and
5wt.%Cr alloys, nitrided at
520°C during 48h with Kn 1.
Thermo-Calc phases
simulation at 50µm from the
surface and at LCth, for each
alloy.
Fig. 1- Profili della
concentrazione totale di azoto
in rapporto alla profondità di
nitrurazione delle leghe Fe -
1, 3 e 5wt.% Cr, nitrurate a
520°C per 48h con Kn 1.
Simulazione delle fasi
mediante Thermo-Calc a
50µm dalla superficie ed alla
LCth(mass nitrogen Limited
Concentration), per ogni lega.
Table 2 – Three binary alloys
nitrided at 520°C during 48h
with a Kn of 1. Determination
of nitrogen excess at 50µm
from the surface, from LCth
(wt.%N) and experimental
mass nitrogen concentration
determined at 50µm from the
surface values.
Tabella 2 – Tre leghe binarie nitrurate a 520°C durante 48h con un Kn di 1. Determinazione dell'eccesso di azoto: a 50µm dalla superficie,
dalla LCth (mass nitrogen Limited Concentration - N peso %), concentrazione totale di azoto determinata sperimentalmente a 50µm dai
valori di superficie.
30
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TRATTAMENTI TERMICI
Fig. 2 – Mass Nitrogen
concentration profiles versus
nitriding depth of Fe -
3wt.%Cr alloy, nitrided at
520°C during 48h with Kn 1,
then annealed at 520°C
during 48h (from electron
microprobe analysis).
Thermo-Calc phases
simulation at 50µm from the
surface after nitriding and at
LCth. Micrograph of the
nitrided zone showing
discontinuous precipitation
below the surface.
Fig. 2 – Profili della
concentrazione totale di
azoto (dall’ analisi con
microsonda elettronica) in
rapporto alla profondità di
nitrurazione nella lega Fe - 3
% in peso Cr, nitrurata a 520°C per 48h con Kn 1, quindi sottoposta a trattamento a 520°C per 48h. Simulazione delle fasi mediante
Thermo-Calc a 50µm dalla superficie ed alla LCth. Micrografia della zona nitrurata che mostra la precipitazione discontinua sotto la
superficie.
Memorie
Fig. 3 (left) – Fe-3wt.%Cr, nitrided at 520°C during 48h, nitrides
extracted on carbon replicas at 50µm from the surface,
investigated with HRTEM and EDX. HRTEM precipitate image
and corresponding Fourier Transformation (calculated diffraction
pattern obtained) confirmed that the precipitate is CrN FCC. The
EDX analysis revealed the presence of iron in the nitride. Blank
analysis, close to the precipitate, allowed to identify peaks due to
the environment (copper from the grid, carbon and silicon from
the carbon deposition), and ensure that no matrix was extracted.
Fig. 3 (sinistra) – Lega Fe-3 % in peso Cr, nitrurata a 520°C per
48h; nitruri estratti con repliche di carbonio a 50µm dalla
superficie, esaminata con HRTEM e EDX. L’immagine HRTEM
del precipitato e la corrispondente trasformata di Fourier
(ottenuta l’immagine di diffrazione calcolata) hanno confermato
che si tratta del precipitato CrN FCC. L'analisi EDX ha rivelato
la presenza di ferro nel nitruro. L'analisi di confronto, vicino al
precipitato, ha permesso di identificare picchi dovuti all'ambiente
(rame dalla griglia, carbonio e silicio dal deposito di carbonio) e
assicura che non è stata estratta anche parte della matrice.
(see Table 2). Fig. 1 and Table 2 show that nitrogen concentration
raises with increase of chromium concentration, associated
with a molar fraction nitrides precipitation predicted
by Thermo-Calc, linked to a hardness increase, as reported
by Hekker [2]. These results show that nitrides are directly
responsible for hardness.
Nitrogen concentration enhances with chromium concentration
increase (as shown in Table 2). According to Thermo-
Calc calculations at 50µm from the surface (Fig.1), no significant
solid solution nitrogen excess occurs, but nitrogen
content from LCth depth up to the surface is explained by the
iron precipitation in complex MN nitrides, isomorph of CrN
FCC (about 30at.%). Indeed the iron presence leads to increase
nitrogen concentration involved in nitrides MN precipitation
and enables nitrogen mass balance to be equilibrated.
Note that the iron atomic radius (0.1241nm) is similar to
the chromium radius (0.1249nm) and FeN, obtained by PVD
(lattice parameter 0.430nm), is isomorph of CrN FCC (lattice
parameter 0.414nm). Iron acts like chromium, allowing nitrides
to keep the same size, morphology, semi-coherence with
the matrix and so possibly the same effect on hardness.
In order to check nitrogen mobility, nitrided Fe-3wt.%Cr alloy
was annealed at nitriding temperature and time (520°C,
48h). Fig. 2 shows nitrogen profiles after nitriding and after
annealing versus nitriding depth, and confirms that nitrogen
diffuses to the core. After annealing, nitrogen concentration
profile presents a plateau, coinciding with LCth. Thermo-
Calc results (fraction phases with chemical composition, at
50µm from the surface for both of them and at LCth depth)
are also plotted on Fig. 2. Thermo-Calc indicates that during
annealing, the phase (Fe x
,Cr y
)N tends to a more stable phase
CrN, leading to loss of iron - about 22at.% - (and corresponding
nitrogen atomic fraction) associated with the molar
fraction nitrides decrease by about 33% (9mol.% to
6.8mol.%), which also should contribute to decrease the
hardness, too.
Microstructural characterizations were performed on nitrides
from carbon replicas and thin foils with High Resolution
Transmission Electron Microscopy and Energy Dispersive
X-Ray. For both cases after nitriding and after annealing, at
50µm from the surface, precipitates with platelet morphology,
present at grain boundaries and in discontinuous precipitation,
were confirmed by High Resolution Image Fourier
Transformation to be chromium nitrides CrN FCC. Then
EDX analysis was performed on them, after verification close
to the precipitate that the ferritic matrix was not extracted.
In the first case, after nitriding, EDX analysis (Fig.3) confirmed
that nitrides (10 to 20nm diameter, 2-10nm thick) contain
iron, with an iron fraction distribution of 14-37at%, in
agreement with Thermo-Calc prediction (22at.%Fe).
In the second case, after annealing, Thermo-Calc predicts an
iron concentration of 8at.% in nitrides. Actually, the coarser
nitrides (20nm diameter), after EDX analysis (Fig.4), contain
from 5 to 14at.%Fe. These investigations confirmed
that nitrides lost iron and corresponding nitrogen (1 to 1
atom) during annealing, the released nitrogen diffusing to
the core.
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TRATTAMENTI TERMICI
7-8/2006 Memorie
Fig. 4 (right) – Fe-3wt.%Cr, nitrided at 520°C during 48h then
annealed at 520°C during 48h, nitrides extracted on carbon
replicas at 50µm from the surface, investigated with HRTEM and
EDX. HRTEM precipitate image and corresponding Fourier
Transformation (calculated diffraction pattern obtained)
confirmed that the precipitate is CrN FCC. The EDX analysis
revealed the presence of iron in the nitride, but less than in
previous case (Fig.3). The EDX analysis revealed the presence of
iron in the nitride. Blank analysis, close to the precipitate, allowed
to identify peaks due to the environment (copper from the grid,
carbon and silicon from the carbon deposition), and ensure that
no matrix was extracted.
Fig. 4 (destra) – Lega Fe-3 % in peso Cr, nitrurata a 520°C per
48h quindi sottoposta a trattamento a 520°C per 48h: nitruri
estratti con repliche di carbonio a 50µm dalla superficie,
esaminati con HRTEM e EDX. L’immagine HRTEM del
precipitato e la corrispondente trasformata di Fourier (ottenuta
l’immagine di diffrazione calcolata) hanno confermato che il
precipitato è CrN FCC. L'analisi EDX ha rivelato la presenza di
ferro nel nitruro, ma in quantità minore rispetto al caso di Fig.3.
L'analisi di confronto, vicino al precipitato, ha permesso di
identificare i picchi dovuti all'ambiente (rame dalla griglia,
carbonio e silicone dal deposito di carbonio) e assicura che non è
stata estratta anche parte della matrice.
II - Industrial steels
A nitriding treatment (520°C 48h Kn3) on 14 different industrial
steels put into evidence the well known preponderant
roles of chromium and aluminium on nitrogen concentration
and hardness profiles. As in binary alloys, a chromium content
increase leads to nitrogen concentration and hardness
enhancements, due to a semi-coherent fine MN FCC precipitation
fraction augmentation. Two steels were selected to
be representative of chromium effect (K 3wt.%Cr) and aluminium
effect (I 1.73wt.%Cr 0.94wt.%Al).
The same methodology applied to binary alloys was used to
study both industrial steels:
• Experiment and first characterization : nitriding then nitrogen
concentration and hardness profiles performed,
• Phases (molar fraction and chemical composition) simulations
by means of the software Thermo-Calc,
• TEM and EDX characterization on nitrides and carbides.
A - Chromium effect: K steel
K steel, which contains 3wt.%Cr, was nitrided at 520°C during
48h. Nitrogen concentration and hardness profiles were
performed then Thermo-Calc calculations at 50µm from the
surface. Hardening (HV x
- HV initial
) is represented as a nitrogen
concentration function (Fig.5). Yet Thermo-Calc predicts
that nitrides fraction keeps on rising linearly with nitrogen
concentration. It can be deduced then that nitrides
precipitation contributes directly to hardening. Another interesting
result of Thermo-Calc calculations (molar fraction
phase versus nitrogen concentration) is that nitride precipitation
requires the entire carbides transformation, in order to
release chromium, vanadium and molybdenum.
Thermo-Calc does not make the difference between vanadium
carbide, vanadium carbonitride and chromium nitride,
as they crystallize in the same structure FCC with a similar
lattice parameter. As reported by Locquet [10], vanadium
carbides, present before nitriding, transform into carbonitride
V(C,N), keeping the same size, morphology and localization.
Considering that result, they should keep the same molar
fraction, which is small (0.58mol.%) compared to the
MN+V(C,N) one, predicted by Thermo-Calc at 50µm from
the surface (9.52mol.%).
At 50µm from the surface, nitrogen mass concentration raises
to 1.27wt.%N. Considering that LCth, calculated including
chromium, vanadium and molybdenum, is 1.00wt.%N,
a “N excess” was revealed equal to 0.28wt.%N. Nitrogen
concentration is once again explained, according to Thermo-
Calc calculations, by the iron alloying in MN complex nitrides
(M stands for Molybdenum, Vanadium and Chromium),
isomorph of CrN FCC. Iron, molybdenum and vanadium
seem to act like chromium and keep the same role towards
hardening mechanisms (especially MN FCC semi-coherent
nitrides).
Some investigations on thin foils at 50µm from the surface
after nitriding, which revealed two different MN precipitations,
as reported by Locquet [10]. Indeed globular nitrides
indexed as MN FCC germinate at the matrix/carbide interface,
due to carbides transformation. Diffraction patterns and
corresponding dark field revealed the presence of semicoherent
finer nitrides MN FCC, with platelet morphology,
Fig. 5 – Hardening (HVx -
HVinitial) representation as a
function of mass nitrogen
concentration of I and K
steels, nitrided at 520°C
during 48h with a Kn of 3.
Phases (molar fraction and
chemical composition)
simulation by Thermo-Calc at
50µm from the surface.
Fig. 5 – Rappresentazione
dell’indurimento (HVx -
HViniziale) in funzione della
concentrazione totale di azoto
negli acciai K e I, nitrurati a
520°C per 48h con un Kn di
3. Simulazioni di fasi
(frazione molare e
composizione chimica)
mediante Thermo-Calc a
50µm dalla superficie.
32
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TRATTAMENTI TERMICI
due to the nitrogen and elements dissolved in the matrix
combination.
Presently, hardening mechanisms could be related as follows
(from [10,11] and Thermo-Calc calculations):
• Fine semi-coherent nitrides MN FCC (M stands for Fe,
Cr, Mo and V) precipitate from the combination of nitrogen
and elements dissolved in the matrix. They are known
to contribute in a large part to nitriding hardening.
• M 23
C 6
carbides (inter- and intragranular) transform into
globular mixed (Fe, Cr, Mo, V)N FCC nitrides. The chemical
composition of these nitrides will be soon investigated.
Their contribution to nitriding hardening is not already
clearly identified.
• Vanadium carbides present before nitriding transform into
carbonitride V(C,N), and can not induce more hardness.
Memorie
B - Aluminium effect: I steel
Both steels I (1.73wt.%Cr and 0.94wt.%Al) and K (3wt.
%Cr) were nitrided at 520°C during 48h with a Kn of 3. Nitrogen
concentration and hardness profiles were performed.
Then the hardening (HV x
- HV initial
) has been represented as
a function of mass nitrogen concentration, Fig. 5. Even
though both steels have the same CLth (the same nitrogen
concentration involved in (Cr, Mo, V, Al)N precipitation),
steel I shows a nitrogen concentration at 50µm from the surface
of 2.6wt.%N (1.3wt.%N for steel K), and hardening
(∆=HV x
-HV initial
) of 763HV (536HV for steel K)!
Nitrogen concentration below the surface is explained first
by thermodynamical simulation and confirmed by HRTEM
and EDX, as the iron alloying in complex MN FCC nitrides
(M = Cr, Mo, V) for both steels I and K. Nevertheless the
iron concentration seems to be limited to about 20at.%.
Thermo-Calc predicts hexagonal AlN precipitation, but they
were not observed. In industrial steels, iron and aluminium
combine with chromium and molybdenum to form (Cr, Mo,
Fe, Al)N FCC precipitates (chemical composition is given
Table 3).
Nitrogen concentration observed in I steel is also explained
below the surface by the γ’ (Fe 4
N) and ε (Fe 2-3
N) precipitation
in the diffusion layer, confirmed by X-Ray analysis, up
to 125µm in depth in the diffusion layer, which corresponds
to 1.8wt.%N.
Presently, hardening mechanisms could be explained as follows:
• Semi-coherent platelets were analyzed and identified as
complex MN (M = Fe, Al, Cr, Mo) FCC. Iron would be
contained in MN from LCth depth up to the surface. These
nitrides, known to contribute mainly to hardening
(∆HV), are the result of the combination of nitrogen and
elements dissolved in the matrix. Thermo-Calc simulation
of the matrix composition before nitriding for steels I and
K (Table 4) shows that the nitrogen concentration involved
in MN semi-coherent precipitation is twice more important
for I steel than for K steel, which could explain the
huge hardening observed below the surface in I steel.
• M 23
C 6
and M 7
C 3
carbides, present in bainitic structure
before nitriding, transform into globular incoherent com-
Fig. 6 (left) – I steel (1.73wt.Cr 0.94wt.%Al), nitrided at 520°C
during 48h. Nitrides were extracted on carbon replicas at 50µm
from the surface, investigated with HRTEM and EDX. HRTEM
precipitate image and corresponding Fourier Transformation
(calculated diffraction pattern obtained) confirmed that elongated
carbide M7C3 transform partially into globular nitrides (Cr, Fe,
Al, Mo)N FCC.
Fig. 6 (sinistra) – Acciaio I (1.73% Cr e 0.94wt.% Al), nitrurato a
520°C per 48h. Sono stati estratti nitruri su repliche di carbonio a
50µm dalla superficie, esaminati mediante HRTEM e EDX.
L'immagine del precipitato HRTEM e la corrispondente
trasformata di Fourier (ottenuta l’immagine di diffrazione
calcolata) hanno confermato che il carburo allungato M7C3 si
trasforma parzialmente in nitruri globulari (Cr, Fe, Al, Mo)N
FCC.
plex MN FCC (M stands for Cr, Al, Mo and Fe - Fe below
the surface up to LCth depth -). Elongated carbides transform
only partially into nitrides (as observed Fig.6), but
globular ones, which are smaller, transform wholly (as observed
Fig.7). These nitrides germinate at the matrix/carbide
interface. Their contribution to nitriding hardening is
not yet clearly identified.
• If γ’ and ε precipitate with well-known needle morphology
in lath interfaces, it can be expected that their contribution
to hardening is not significant.
CONCLUSION
Thermo-Calc calculation associated with microstructural
analysis (HRTEM and EDX) allows the determination of nitrogen
localization, which leads to a better understanding of
chromium and aluminium effects on hardening mechanisms.
Binary Alloys
1- Chromium concentration increase in binary alloys induces
nitrogen content enhancement, linked to MN FCC nitrides
fraction and hardness augmentation, showing that
semi-coherent MN nitrides are responsible for hardening.
Table 3 – I steel nitriding at
520°C during 48h. Results of
EDX analysis (average and
range) performed on nitrides
extracted at 50µm from the
surface, with corresponding
sizes.
I T1 50µm surface M Chemical composition (at.%) Size
MN FCC Al Cr Fe Mo (nm)
Semi-coherent Average / 3 zones 45 18 23 13 3-13
platelets Range 41-49 13-24 19-29 6-17 length
Globular germinated Average / 22 particles 13 73,5 7 7 6-28
at the carbides/α interface Range 3-23 61-82 2-13 1-12 length
Tabella 3 – Nitrurazione
dell'acciaio I a 520°C per
48h. Risultati dell’analisi
EDX (media e intervallo dei
valori) effettuata sui nitruri estratti a 50µm dalla superficie, con le corrispondenti dimensioni.
Bainitic lath interface Average / 4 zones 61 18 8 13 3-9
Range 57-66 12-26 4-12 9-16 length
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TRATTAMENTI TERMICI
7-8/2006 Memorie
Phases molar fraction /
Chemical mass composition α
wt.%N
wt.%Cr α
wt.%Mo α
wt.%V α
wt.%Al MN
α from α
K core 650°C
5.54mol.%M 23
C 6
1.16 0.12 0.02 0.01 0.34
0.08mol.%M7C3, 93.77mol.%α
I core 620°C
2.61mol.%M 23
C 6
, 0.28 0.03 0.98 0.59
2.86mol.%M 7
C 3
, 94.53mol.%α
2- After nitriding, complex MN (M = Fe and Cr), isomorph
of CrN FCC, are observed with a maximum of iron atomic
fraction below the surface, which explains nitrogen
content. Iron acts like chromium, inducing nitrides fraction
and hardening increase. It was confirmed that in the
nitrided Fe-3wt.%Cr alloy, at 50µm from the surface,
nitrides MN FCC, with platelet morphology contain
between 14 and 37at.%Fe (22at.%Fe simulated by Thermo-Calc).
3- After annealing, under temperature effect nitrogen diffuses
to the core. From the surface up to LCth depth, nitrides
MN loose partially iron and nitrogen corresponding
(1 N atom to 1 Fe atom), decreasing MN molar fraction.
From LCth depth up to the core, as all chromium was
not used to precipitate, the released nitrogen can combine
with chromium, to form CrN and so increase CrN
molar fraction and hardening.
Industrial steels
1- After nitriding several industrial a-alloyed steels, it
could have been concluded that chromium and aluminium
enhance nitrogen concentration and hardness profiles,
molybdenum and vanadium having a minor importance.
2- In industrial steels without aluminium, nitrogen mass
balance can be explained, according to Thermo-Calc simulations,
by complex MN nitrides, isomorph of CrN
FCC, containing chromium, iron, molybdenum and vanadium.
Before nitriding, the largest part of chromium,
molybdenum and vanadium precipitated as carbides VC
and M 23
C 6
. These last ones transform during nitriding
into globular incoherent MN FCC. A small fraction of
chromium, molybdenum and vanadium, dissolved in the
matrix, combines with nitrogen and precipitates as semicoherent
platelets MN FCC. More HRTEM and EDX
investigations will be performed to identify chemical
compositions of both kinds of precipitates.
3- In industrial steels, aluminium allows a huge nitrogen
concentration and induces significantly important hardening.
Below the surface, nitrogen mass balance is first
explained, according to Thermo-Calc simulation and
confirmed by characterization, by complex nitrides MN,
isomorph of CrN FCC, containing iron (atomic fraction
seems to be limited to 30at.%), aluminium, chromium
and molybdenum. γ’ and ε phases precipitation in the
diffusion layer allows nitrogen balance to be equilibrated.
Another new important result is that aluminium precipitates
in complex MN nitrides FCC, and not as hexagonal
AlN and contributes to an important hardness increase.
REFERENCES
Table 4 – From Thermo-Calc
simulation of the ferritic
phase chemical composition,
determination of the mass
nitrogen concentration, which
could precipitate as semicoherent
MN FCC nitrides
with elements dissolved in the
matrix - K and I steels -
Simulation at the last pretreatment
before nitriding
(650°C for K and 620°C for I).
Tabella 4 – Da simulazione Thermo-Calc della composizione chimica della fase ferritica, determinazione della concentrazione totale di
azoto, che può precipitare in forma di nitruri semi-coerenti di MN FCC, con M costituito dagli elementi dissolti nella matrice ( acciai I e K ).
Simulazione all'ultimo pretrattamento prima della nitrurazione (650°C per K e 620°C per I).
Fig. 7 (right) – HRTEM precipitate image (and Fourier Transform
not shown here) and EDX analysis confirmed that globular
carbides transform wholly into globular nitrides MN FCC, that
semi-coherent platelets are MN FCC and that precipitates on lath
interfaces are also MN FCC. M stands for iron, aluminium (in
chromium presence, aluminium does not precipitate as hexagonal
AlN), chromium and molybdenum.
Fig. 7 (destra) – Immagine HRTEM del precipitato (la
trasformata di Fourier non è qui riportata) e analisi EDX hanno
confermato che i carburi globulari si trasformano interamente in
nitruri globulari MN FCC, che le piastrine semi-coerenti sono MN
FCC e che anche i precipitati sulle interfacce sono MN FCC. M
sta per ferro, alluminio (in presenza di cromo, l'alluminio non
precipita come AlN esagonale), cromo e molibdeno.
[1] D.H. JACK, K.H. JACK, Materials Science and Engineering
11, (1973) p. 1-27
[2] P.M. HEKKER, E.J. MITTEMEIJER, H.C.F. ROZEN-
DAAL, The influence of Nitriding on the Microstructure
and Stress State of Iron and Steel (1985) p. 51-61
[3] E.J. MITTEMEIJER, H.C.F. ROZENDAAL, P.J. VAN
DER SCHAAF, R.T. FURNEE, The influence of Nitriding
on the Microstructure and Stress State of Iron and
Steel (1985) p. 109-117
[4] B.J. LIGHTFOOT, D.H. JACK, Heat Treatment’73,
London, (1973) p. 59-65
[5] H.J. SPIES, S. BÖHMER, HTM 39 (1984) p. 1-6
[6] Y.M. LAKHTIN, Y.D. KOGAN, Mashinostroenie, Moscow,
(1976)
[7] V.A. PHILLIPS, A.U. SEYBOLT, Trans. Of the Metall.
Society of AIME 242, (1968) p. 2415-2422
[8] H.C.F. ROZENDAAL, E.J. MITTEMEIJER, P.F. CO-
LIJN, P.J. VAN DER SCHAAF, Metall. Trans. A 14, 2
(1983) p. 395-399
[9] J.P. CALVEL, Détermination expérimentale des contraintes
résiduelles introduites par la nitruration gazeuse
d’aciers 35CD4 et 40CAD6.12 - Relations avec le
34
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TRATTAMENTI TERMICI
durcissement, PhD Thesis INPT, France (1983)
[10] J.N. LOCQUET, R. SOTO, L. BARRALLIER, A.
CHARAÏ, Microsc. Microanal. Microstruct. 8 (1997)
p. 335-352
[11] M. SENNOUR, Apport de la microscopie électronique
en transmission et de la spectroscopie EELS à la caractérisation
de nitrures (AlN, CrN) dans le fer et l'alliage
Fe-Cr, PhD Thesis INSA Lyon, France (2002)
[12] E.J. MITTEMEIJER, Journal of Metals (1985) p.16-20
[13] E.J. MITTEMEIJER, M.A.J. SOMERS, Proc. Int.
Conf. on High Nitrogen Steels, Lille, France (1989)
[14] R.E. SCHACHERL, P.C.J. GRAAT, E.J. MITTE-
MEIJER, Metallurgical and Materials Transactions A
35 (2004) p. 3387-98
[15] M.SENNOUR, P.H. JOUNEAU, C. ESNOUF, Journal
of Materials Science 39 (2004) p.1-11
[16] B. MORTIMER, P. GRIEVESON, K.H. JACK, Scandinavian
Journal of Metallurgy 1, (1972) p. 203-209
[17] P.C. VAN WIGGEN, H.C.F. ROZENDAAL, E.J. MIT-
TEMEIJER, Journal of Materials Science (1985) p.
4562-4582
[18] M.A.J. SOMERS, R.M. LANKREIJER, E.J. MITTE-
MEIJER, Philosophical Magazine A 59 (1989) p. 353-
378
[19] R.E. SCHACHERL, P.C.J. GRAAT, E.J. MITTE-
MEIJER, Z.Metallkunde 93 (2002) 5 p.468-477
[20] E.J. MITTEMEIJER, M.A.J. SOMERS, Surface Engineering
13 (1997) 6
Memorie
ABSTRACT
UN NUOV0 APPROCCIO AI MECCANISMI DI INDURIMENTO
NELLO STRATO DI DIFFUSIONE DEGLI ACCIAI LEGATI,
CON STRUTTURA α NITRURATI IN FASE GASSOSA
EFFETTI DI CROMO E ALLUMINIO:
STUDIO SPERIMENTALE E DI SIMULAZIONE
Parole chiave: trattamenti termici, nitrurazione, acciaio,
metallurgia fisica, caratterizzazione materiali
Sono stati studiati i meccanismi di indurimento nello strato
di diffusione di ferro e acciai con struttura α nitrurati in fase
gassosa, attraverso l’analisi degli effetti della presenza di
cromo (leghe binarie ed acciai industriali) e di alluminio
(acciaio industriale). Dopo la nitrurazione (520°C 48h), il
bilancio di massa dell’azoto fra la concentrazione di azoto
totale nella zona di diffusione, determinata sperimentalmente,
e la concentrazione di azoto prevedibile in via teorica, rivela
per ogni lega "un eccesso di azoto". Jack e Mittemeijer
[1-3 ] hanno suggerito che le differenze in volume fra i nitruri
semi-coerenti e la matrice induca la distorsione locale
del reticolo della matrice, e ciò porti ad un aumento locale
della solubilità dell'azoto nella matrice. Nel presente studio
si propone un nuovo approccio, basato su calcoli termodinamici
(software Thermo-Calc), confermati mediante diversi
metodi di caratterizzazione (HRTEM, EDX e raggi X).
In effetti non si riscontra alcuna significativo "eccesso di N"
in soluzione solida, ma la concentrazione totale nell'azoto si
spiega con la precipitazione di nitruri complessi metalloazoto
(MN), isomorfi di CrN cubico a facce centrate, contenenti
cromo, ferro (fino a 30 % atomico, a 50µm dalla superficie),
molibdeno e vanadio. Durante il trattamento
(520°C 48h), la frazione atomica di ferro nei nitruri complessi
(MN) diminuisce e la corrispondente frazione atomica
di azoto si diffonde all’interno.
L’aggiunta di alluminio nell’acciaio industriale permette di
aumentare decisamente la concentrazione di azoto e l’indurimento
(∆=HV x
-HV iniziale
). L'alluminio induce, nello strato
di diffusione, precipitazione di Fe 4
N e di Fe 2-3
N e di precipitati
nei nitruri complessi metallo-azoto (MN) a struttura
FCC, contenenti cromo, ferro e molibdeno.
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TRATTAMENTI TERMICI
The behaviour of decarburized layers
in steel nitriding
I. Calliari, M. Dabalà, E. Ramous, M. Zanesco, E. Gianotti
Samples of quenched and tempered 40CrMo4 steel, previously surface decarburized at different depths,
have been submitted to gaseous nitriding. The surface layers after decarburization and nitriding was
examined by optical microscopy (OM), scanning electron microscopy (SEM) and microhardness tests.
The distribution of iron nitrides in the diffusion layer was analyzed by X-ray diffraction (XRD). The
nitrogen and carbon profiles in the diffusion layers were determined by scanning electron microscope
equipped with a wavelength dispersive spectrometer (EPMA-WDS).
For all the specimens, the highest hardness values only change slightly with increasing time and
temperature of decarburization, while case hardness depths decrease.
In all the specimens, the nitriding depth, as determined by the nitrogen profile, is larger than the one
determined by the hardness profile.
Memorie
Key words: decarburization, nitriding, steel, 40CrMo4
INTRODUCTION
The surface microstructure on nitriding steels has been widely
investigated and discussed in several papers (1-4). The
nitriding process develops a zone which can be subdivided
into a compound (white) layer, with carbo-nitrides, and an
underlying diffusion layer. The compound layer is formed
by nitrides and carbonitrides, while interstitial nitrogen and
alloy forming elements (Cr, Mo, Al, and V) nitrides are present
in the diffusion layer.
The nitrided layers structure depends on the interaction
between the diffusional nitrogen and original steel microstructure,
normally ferrite and carbides , caused by the tempering
heat treatment, which is usually performed on steel
before nitriding. A special role is played by the transformation
of Cr carbides to Cr nitrides (5), so the C may diffuse
interstitially from the diffusion layer to the compound layer,
where it forms carbides. This transformation explains why
the final microstructure of nitrided steels depends not only
on the C content but also on the C distribution(6) in the prior
steel microstructure. The nitriding treatment modifies the
carbides and C distribution which then affects also the hardness
profile in the N-enriched layer(7). Occasionally, in industrial
practice, the nitrided work pieces can present some
surface decarburization due to previous heat treatments. The
behaviour of nitrided samples after a controlled surface decarburization
has been examined in this study.
The main differences between decarburized and non-decarburized
layers are obviously the lower carbides content and
the larger ferrite grain sizes in the former. The low carbide
content determines that nitrides may form by precipitation
from ferrite and not by carbides decomposition. The large
ferrite grain size implies that the grain boundary nitrogen
diffusion increases and the grain boundary nitrides precipitation
is enhanced.
Moreover, the study could give more information about the
I. Calliari, M. Dabalà, E. Ramous, M. Zanesco
DIMEG Department of Innovation in Mechanics and Management, Padova
E. Gianotti
Heat Treatments Ferioli & Gianotti Spa, Rivoli (TO)
Paper presented at the (2nd International Conference
HEAT TREATMENT AND SURFACE ENGINEERING IN AUTOMOTIVE APPLICATIONS,
organised by AIM, Riva del Garda, 20-22 June 2005
effect of pre-oxidation treatments, often suggested to improve
nitrogen take-up, and about the redistribution of carbon
observed during nitriding.
EXPERIMENTAL PROCEDURES
The specimens have been obtained from a bar (diameter 28
mm) of commercial steel 4140 quenched (in oil at 840°C)
and tempered (2 h at 600°C).
Six different thermal cycles, aimed at obtaining surface decarburization,
were conducted in the experimental condition
reported in table 2.
After metallographic preparation, the cross-sections of the
specimens were etched with Nital 4% in order to characterize
the structure and evaluate the decarburized thickness through
optical microscopy.
The decarburized specimens were submitted to the following
nitriding treatment: all the specimens were heated at
350°C in air, from 350°C to 400°C in N 2
and then kept at
400°C for 30 minutes in N 2
, from 400°C to 510°C in N 2
and
then kept at 510°C for 72 h with a NH 3
dissociation of 25%.
Steel
Composition %
C Mo Mn Cr Ni Cu Si
4140 0,432 0,2 0,85 1,148 0,122 0,267 0,243
Table 1 – Composition of the nitrided steel.
Tabella 1 – Composizione dell’ acciaio.
Sample Thermal cycle Cooling
1 3 h 800°C Air
2 4 h 800°C Air
3 3 h 840°C Air
4 4 h 840°C Air
5 3 h 900°C Air
6 4 h 900°C Air
Table 2 – Decarburization heat treatment.
Tabella 2 – Trattamenti per ottenere gli strati decarburati.
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TRATTAMENTI TERMICI
7-8/2006 Memorie
The cross-sections of the nitrided specimens were etched
with Nital 4% to observe the microstructure and subsequently
etched with Murakami in order to evaluate the distribution
of carbides, by light and SEM microscopy. Microhardness
profiles were performed using a load of 200g. The concentration
profiles of nitrogen and carbon in the nitride
layers have been determined by GDOS spectrometry on the
surface layers, and by the JEOL-JXA microprobe (with
spectrometers 4 WDS). A Siemens D5000 X-ray diffractometer
with Cu Kα radiation was used to examine the phases
of white surface layers.
RESULTS AND DISCUSSION
1. Decarburized samples
Table 3 reports the values of decarburized depths, surface
hardness after nitriding and effective depth of nitriding ; it
is evident that the decarburized thickness increases with time
and temperature of heat treatment.
In fig.1 an example(sample 2, 800°C,4h) of the microstructure
of the decarburized surface layer is reported. Particularly
in the samples n. 5 and 6, the decarburization produces
a completely ferritic
microstructure with coarse grains. The increase in ferrite
grain size is important because, during the nitriding process,
the nitrogen diffuses preferentially along the grain bounda-
Surface Effective
Depth of hardness depth
Sample Decarbu- after after
rising (µm) nitriding nitriding
(HV 0.5
) (µm)
0 0 841+-12 481
1 150 817+-10 476
2 290 836+-16 470
3 300 810+-12 444
4 310 841+-12 415
5 460 814+-14 436
6 490 817+-16 388
Table 3 – Decarburized depths, surface hardness after nitriding
and effective depth of nitriding.
Tabella 3 – Spessori degli strati decarburati, durezze superficiali
dopo nitrurazione e spessori efficaci di nitrurazione.
Fig. 1 – Optical micrograph of sample 2 (800°C,4h).
Fig. 1 – Microstruttura del campione 2 dopo decarburazione
(800°C, 4 ore).
Fig. 2 – Hardness profiles of samples 5,6 and 0.
Fig. 2 – Profili di microdurezza dei campioni n. 5, 6 e 0.
ries. As the grain size increases, the nitrogen gradient concentration
between grain core and grain boundaries may increase
and the nitrides precipitation along the grain boundaries
is favoured.
2. Nitrided samples
2a. Surface hardness
The surface hardness values measured after nitriding are reported
in the 3 rd column of table 3. The values are the mean
of 10 measurements randomly taken on the surface. Samples
from 1 to 6 have been heat treated as indicated in table
2, while sample 0 has only been nitrided in the same experimental
conditions of 1-6 but not decarburized.
All the hardness values are higher than the 600HV required
by the standard. The surface hardness is not affected by the
different rate of decarburization, as sample 0 has the same
hardness as the others. High nitrogen enrichment in the
compound layer probably compensates the carbon loss and
the resulting hardness can only be determined by nitriding.
2b. Hardness profiles
Fig. 2, reports the microhardness profiles of samples 5 and
6, superimposed with profiles of sample 0. All curves have
the typical trend of nitrided steels, showing a decreasing
hardness from surface to core. There is no sharp transition in
the hardness values from the compound layer to the diffusion
zone and this can be attributed to the high hardness levels
achieved in the diffusion zone after the long nitriding
process and despite the carbon loss. The effect of decarburization
can be evaluated by comparing the profiles of samples
1-6 with the profile of sample 0 in the diffusion zone.
The hardness values of sample 0 are slightly higher than
those of samples 1-6; the differences are in the range of 10-
20HV and decrease as the depth increases. This difference
can be explained by the lower hardness values of the decarburized
samples where a limited amount of carbides is present.
The consequence of the lower profile is that the effective
case depth after nitriding decreases. The last column of
table 3 reports the effective depth values measured on all
the samples, the hardness of 460HV has been achieved at
depths varying from 390 to 480microns. The value of
480microns has been measured in sample 0, while sample 6
(having the deepest decarburization) has the minimum depth
of 390 microns.
2c. Microstructure
Optical and electronic microscopy metallographic investigations
were performed to observe the influence of decarburization
on the microstructure after nitriding. The results for
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TRATTAMENTI TERMICI
Memorie
Fig. 3a – SEM micrograph of sample 0.
Fig. 3a – Micrografia SEM del campione 0.
Fig. 3b – SEM micrograph of sample 5.
Fig. 3b – Micrografia SEM del campione 5.
sample 0 (nitriding without decarburising) and sample 5
(850°C, 4h and nitriding) are presented in figs. 3a and 3b.
The compound layers and diffusion zone with grain boundary
precipitation are visible.
The main features of the white layer in the decarburized
samples can be summarized as follows:
- It is about 30% thinner than in simply nitrided sample,
- It appears unhomogeneous, with many voids, diffuse porosity
and some cracks and the outer surface is friable
with scaling, mainly in the deeply decarburized samples.
In the diffusion zone, dark precipitates formed at the grain
boundaries. Fig.4 reports the SEM-BSE micrograph taken
without etching on sample 3 cross section (800°C, 30’).The
precipitation starts beneath the white layer and grows for
about 10 microns towards the diffusion layer. This morphology
is more evident in highly decarburized samples, where
the high temperature and holding time of heat treatment determines
grain growth. After this heat treatment the surface
base material was partially ferritic and the ferrite volume
fraction increased with temperature and holding time. After
nitriding, no significative amount of ferrite was detected in
the low decarburised samples. This means that interstitial
diffusion of N and C occurs, both at grain boundaries and inside
the grains, resulting in a quite homogeneous microstructure.
On the other hand, in the diffusion zone of highly decarburised
samples some ferritic grains still remain. It means that in
these samples the diffusion mechanism occurs mainly at
grain boundaries and despite the long nitriding time, the interstitial
elements C and N were not able to diffuse up to the
inner zone of the ferritic grains.
2d. Composition and microstructure of nitrided layers
The concentration profiles of carbon and nitrogen in crosssections
of decarburised and nitrided samples were determined
by wave- length dispersive electron microprobe analysis
(WDS-EPMA).
The results, obtained from samples 0 and 5 (900°C, 3h, nitriding),
are given in figs, 5a and 5b, along with the hardness
profiles for comparison. The C and N concentrations are expressed
as Count/sec/mA.
The higher nitrogen concentrations, measured on the surface
of sample 0, are caused by the compound layer, which is less
porous and more homogeneous than in samples 3 and 5.
In all the samples the nitrogen content decreases from surface
to core; with the mean concentration in the diffusion zones
of samples 3 and 5 slightly lower than in sample 0. In
all specimens the nitriding depth as determined by the nitrogen
profile is larger than that determined by the hardness
Fig 4 – SEM-BSE micrograph of sample 3.
Fig. 4 – Micrografia SEM-BSE del campione 3.
profiles, arriving at about 700microns.
The carbon profiles are obviously more affected by the decarburising
process. The peaks in the carbon concentrations
in the diffusion layer of sample 0 indicate the presence of
carbides. The carbon profiles in the diffusion zone of samples
3 and 5 show the different depth of decarburised layer.
Some peaks in profile of sample 3 indicate the presence of a
few (small) carbides and the profile of sample 5 indicates a
complete absence of C beneath the surface. The carbon concentration
profiles of both samples correspond to the depth
of the decarburized layers evidenced by the optical microscopy.
The carbon and nitrogen contents in compound layers have
been examined by the Glow Discharge Optical Spectrometry
(GDOS), to a depth of about 70 microns with nanometric
spatial resolution.. The profiles of C and N obtained on samples
0 and 6 are reported in fig.6 The GDOS profile of carbon
in sample 0 after nitriding show that carbon is absent for
about 10 microns, which correspond to the white layer, and
increases with depth to reach the base alloy concentration of
0.40%.The carbon loss associated with nitriding is well
known: during the nitriding treatment the carbon diffuses to
the core as well as to the surface.
The GDOS profiles of C and N in decarburized samples show
that carbon is absent in the analysed depth (about 70 microns).
Instead, there is a higher nitrogen surface concentration (more
than 1%), compared to to that observed in the sample without
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TRATTAMENTI TERMICI
7-8/2006 Memorie
Fig. 5a (left) and fig. 5b (right) – Nitrogen and Carbon EPMA-WDS profiles and hardness profile of sample 0 (left) and of sample 5 (right).
Fig. 5a (sinistra) e 5b (destra) – Profili di microdurezza e delle concentrazioni di carbonio e azoto ottenuti con EPMA-WDS del campione 0
(a sinistra) e 5 (a destra).
Fig.6 – GDOS profiles of N and C taken on the compound layers of sample 0 (left) and 6 (right).
Fig. 6 – Profili di concentrazione del carbonio e dell’ azoto nello strato di composti, ottenuti con GDOS nei campioni 0 (a sinistra) e 6 (a
destra).
decarburization, followed by a smooth decrease to a value
which remains constant in the examined layer.
This higher nitrogen concentration in the surface layers,
both in the white layer and beneath this layer, appears as the
main difference, when compared with the non-decarburized
sample. The nitrogen content of the white layer in the sample
0 is about 0,4%, and arrives to the value of 1, 5% in the
more decarburised samples. There is a similar difference in
the inner layer, beneath the compound one. These data suggest
that the surface decarburisation favours the diffusion of
nitrogen, mainly along the grain boundaries
These results have been confirmed by the metallographic investigation,
which evidenced that in decarburised samples the
nitrogen compounds precipitate at grain boundaries as needlelike
particles, coarser in the more carbon depleted samples.
The GDOS measurements evidenced also the surface enrichment
in Si, Mo and Cr in the decarburised and nitrided
samples. The GDOS profiles indicate that the Si, Mo and Cr
enrichment is limited to the 5-10 microns of the white layer.
The Si and Mo concentrations measured at the surface are
about 2 %, while the Cr concentration reaches 7-9%. It has
been verified that surface enrichment occurs during the decarburisation
heat treatment and is not affected by nitriding.
Fig.7 reports the Si, Mo and Cr profiles obtained on the decarburized,
not nitrided, sample 5.
Fig.7 – GDOS profiles of Mo,Si,Cr on sample 5.
Fig. 7 – Profili di concentrazione di Mo, Si e Cr ottenuti con
GDOS nel campione 5.
The occurrence of Si, Mo and Cr surface enrichment can be
easily explained by their high oxygen affinity which favours
their outward diffusion during the air decarburisation treatment.
18
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TRATTAMENTI TERMICI
The surface layers constitution was investigated with XRD:
in the white layer the presence of nitrides ε (Fe 3
N) e γ
(Fe 4
N) was confirmed. The volume fraction of ε is about
60% in all samples; a moderate increase of γ fraction was
detected in the more deeply decarburized samples.
CONCLUSIONS
The effects of decarburisation on properties and structure of
nitrided surface layers in the 40Cr4Mo steel can be summarized
as follows:
1. The surface hardness is not affected and the hardness
profiles of decarburised samples are slightly lower than
the profile of the only nitrided sample.
2. In all samples the nitrogen content decreases from surface
to the bulk. In the previously decarburised samples
the nitrogen concentration is significantly higher in the
first surface layers and slightly lower in the diffusion zones,
as compared with the non-decarburised sample.
3. In all specimens the nitriding depth as determined from
the nitrogen profile is larger than that determined from
the hardness profiles, arriving at about 700 microns.
4. The white layer of nitrided and decarburised samples is
more porous and friable, with coarser grains than in the
non-decarburized samples.
5. The thickness of the white layer decreases as the decarburisation
increases.
6. The main effect of the previous decarburization on the
nitrided layers seems to be the modification of the
morphology and composition of the compound layer.
7. The actual effects of carbon loss on microstructure and
hardness after nitriding are probably compensate by the
strong nitriding conditions of our experiment, which favour
the growth and development of compound layers.
REFERENCES
(1) T.Hirsch, F.Hoffmann, P.Mayr, Surf.Eng., 14, (1998), p.
481
(2) R.S.Schneider, H.Hbler, J.Mat.Sci., 33, (1998) p. 1737
(3) B.M. Korevaar, S. Coorens, Y. Fu, J. Sietsma, S. van der
Zwaag, Mat. Sci. Techn., 17,
54-62 (2001)
(4) J.Ratajski, J.Tacikowski, M.A.J.Somers, Surf.Eng., 19,
(2003), p. 285
(5) P.C. Van Wiggen, H.C.F. Rozendaal, E.J. Mittemeijer, J.
Mat. Sci., 20, (1985) p. 4561
(6) P. Egert, A.M. Maliska, H.R.T. Silva, C.V. Speller, Surf.
Coat. Techn., 221, (1999), p.33
(7) P. Filippi, G. Rossetti, Met. Ital., 67, (1975) p. 605
Memorie
ABSTRACT
COMPORTAMENTO DI STRATI SUPERFICIALI DECARBURATI
SOTTOPOSTI A NITRURAZIONE
Parole chiave: decarburazione, nitrurazione, acciaio 40CrMo4
Normalmente il trattamento di nitrurazione viene eseguito su
acciai bonificati, quindi con una struttura costituita da martensite
rinvenuta e carburi. L’arricchimento in azoto per diffusione
prodotto dalla nitrurazione modifica questa microstruttura:
in particolare i carburi di cromo vengono trasformati in
azoturi, liberando il carbonio che diffonde sia nella zona di
diffusione, sia verso lo strato superficiale di composti. Quindi
la nitrurazione non solo arricchisce lo strato superficiale in
azoto, ma modifica anche la distribuzione del carbonio: tutto
questo influisce sul profilo e sui valori di durezza degli strati
nitrurati. Poiché nella pratica i trattamenti eseguiti prima della
nitrurazione occasionalmente possono provocare una certa
decarburazione nella zona superficiale dei pezzi, è stato esaminato
il comportamento durante la nitrurazione di strati superficiali
volutamente decarburati in un tipico acciaio da bonifica,
l’AISI 4140 (tabella 1). La decarburazione è stata ottenuta
mediante riscaldamento in aria a temperature comprese
fra 800 e 900°C (tabella 2). Gli strati decarburati (figura 1),
per spessori compresi fra 150 e 300 micron, presentavano una
struttura con grossi grani di ferrite ed ovviamente assenza di
carburi. La nitrurazione è stata eseguita a 512°C per 72 ore,
con ammoniaca dissociata al 25%. I campioni trattati sono
stati esaminati con le usuali tecniche metallografiche, spettrometria
GDOS, microsonda WDS e diffrazione X. Dopo nitrurazione,
la durezza superficiale dei campioni decarburati è solo
lievemente inferiore a quella del campione solo bonificato,
e comunque largamente superiore al valore minimo richiesto
(tabella 3). Anche i profili di microdurezza (figura 2), mostrano
solo una lieve diminuzione rispetto ai valori riscontrati sul
campione non decarburato. Però questa diminuzione si mantiene
per tutta lo spessore dello strato trattato e di conseguenza
i campioni decarburati presentano una profondità efficace
di nitrurazione inferiore (tabella 3). Tale diminuzione aumenta
con l’ aumentare della profondità di decarburazione. Questo
effetto è certamente dovuto alla minore concentrazione di carbonio
negli strati decarburati. Nella microstruttura degli strati
nitrurati dopo decarburazione (figura 3 a,b e c) le maggiori
differenze riguardano lo strato di composti, che ha uno spessore
inferiore di circa il 30%. Inoltre questo strato appare eterogeneo,
con molti vuoti, porosità e fessurazioni, che lo rendono
friabile e poco resistente allo scagliamento. Nello strato di
diffusione si notano evidenti fenomeni di precipitazione a bordo
grano, che confermano la prevalente diffusione dell’ azoto
lungo i bordi grano. Tuttavia grani di ferrite sono presenti solo
nei campioni maggiormente decarburati. Con analisi
WDS-EDS sono stati tracciati i profili di concentrazione del
carbonio e dell’azoto nello strato di diffusione (figura 5 a e b).
Nei profili del carbonio si notano i picchi dovuti ai carburi nel
campione non decarburato, mentre il carbonio è praticamente
assente nel campione decarburato. Invece in tutti i campioni l’
andamento del profilo dell’ azoto presenta una diminuzione
graduale, che arriva a profondità superiori a quelle corrispondenti
all’ aumento della durezza, a conferma che piccole
concentrazioni di azoto possono avere solo scarsa influenza
sulla durezza. Il contenuto di carbonio e azoto nello strato superficiale
di composti è stato invece analizzato con la spettrometria
GDOS (figura 7 a e b). Per il carbonio i risultati mettono
in evidenza la ovvia assenza di carbonio nello strato superficiale
dei campioni decarburati, mentre nel campione solo nitrurato,
il carbonio è assente solo fino a circa 10 micron, a
conferma dell’ effetto di decarburazione superficiale prodotto
dalla nitrurazione. L’ andamento dell’ azoto evidenzia invece
che lo strato di composti nei campioni decarburati ha una
concentrazione di azoto sensibilmente più elevata, circa 1,5%,
rispetto al campione solo nitrurato, dove l’ azoto arriva allo
0,4%. Nei campioni decarburati anche lo strato di diffusione
più interno, presenta un’ analoga maggiore concentrazione di
azoto. Ciò sembra indicare che la decarburazione favorisce
una maggiore diffusione dell’azoto, probabilmente soprattutto
lungo i bordi grano. Le analisi GDOS hanno anche rivelato un
arricchimento superficiale in molibdeno, silicio e cromo (figura
7), tutti elementi in lega con elevata affinità per l’ossigeno.
Tale arricchimento è giustificato da una diffusione preferenziale
di questi elementi verso la superficie, durante il trattamento
di decarburazione ossidante.
7-8/2006
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TRATTAMENTI TERMICI
Surface hardening of Al 7075 alloy
by diffusion treatments
of electrolytic Ni coatings
K. Brunelli, M. Dabalà, C. Martini
A surface hardening process based on Ni coating and subsequent diffusion heat treatments was studied
for Al7075 alloy. Nickel coatings with different thickness on 7075 Al alloy were obtained by electrolytic
and electroless processes . Heat treatments in inert atmosphere at 500 °C and 530 °C for different times
were performed in order to obtain surface hardening of the aluminum alloy by diffusion of Ni into the
substrate. The effect of temperature and time on the depth of hardening was studied by SEM, EDS, Glow
discharge optical spectrometry, microhardness tests and tribological tests.
Surface hardness higher than 1000 HV and depth of hardening higher than 100 mm were achieved by
diffusion of Ni layers and formation of Al 3
Ni 2
and Al 3
Ni intermetallic phases .
The tribological test on Al7075 alloy against a hard Cr coated steel cylinder in air showed an average
coefficient of friction µ of about 0.5 with a wide variation range, while all the coated and treated samples
exhibited a coefficient friction of about 0.7. However, the wear scar depth of the hardened alloy is about
30 times lower than that of the Al 7075 alloy.
Memorie
Key words: Aluminum alloys, Hardening, Ni coating, Heat treatments, Electron Microscopy, Wear
INTRODUCTION
Aluminum and its alloys are attractive for many application
in chemical, automobile and aerospace industries because of
their excellent properties as height strength-to- weight ratio,
high electrical and thermal conductivities and good formability.
However their hardness, wear resistance and mechanical
properties are poor in comparison to steel resistance and
continuous efforts are made in the research into new possibilities
for making use of the advantage of the aluminum in
application that were reserved up to now for harder and more
wear-resistant materials. The solution mainly adopted is
to produce a thick hardened layer on the substrate by laser
surface alloying process [1-2], plasma vapor deposited PVD
[3], chemical vapor deposited CVD [4] and diamond-like
hydrocarbons coatings [5].
Recently, a method based on Ni-B electroless plating followed
by diffusion heat treatment was proposed for improve
the wear resistance of titanium alloys [6]. The diffusion of
nickel and boron present in the coating allowed a remarkable
hardening of the titanium alloy. It has been chosen to
use the same method for the aluminum alloys. In this work,
the effect of the diffusion of nickel and boron coating, obtained
by electroless deposition, and Ni coating, obtaining by
electrolytic process, on the hardening of the Al 7075 was
examined.
EXPERIMENTAL
Specimens of 7075 Al alloy with a surface area of 12 cm 2
were obtained from a untreated rod and mechanically polished
using standard metallographic procedures. The surface
Katya Brunelli, Manuele Dabalà
DIMEG – Università di Padova - Via Marzolo, 9 35131 – Padova
Carla Martini
Istituto di Metallurgia- Università di Bologna - V.le Risorgimento, 4- Bologna
Component
Electroless
bath
Electrolytic
bath
Nickel chloride 24 g l -1 -
Sodium acetate 36 g l -1 -
DMAB 10 g l -1 -
Sodium lauryl sulphate 0.1 g l -1 -
Nickel sulphate hexahydrate - 225- 410 g l -1
Nickel(II)-chlorid-hexahydrat - 30-60 g l -1
Boric acid - 30-40 g l -1
pH 7 2-3
Temperature 65-70 °C 55 °C
Voltage - 1 V
Current Density - 0.1-0.2 A/cm 2
Table 1 – Composition of the deposition baths.
Tabella 1 – Composizione dei bagni di deposizione.
of the specimens was degreased with alcohol and air dried,
then activated by chemical etching in a 6 % HF aqueous solution
for 5 seconds. After chemical etching, the specimens
were rinsed with a deionized water and immersed either in
the solution for the electroless Ni-B deposition or in the solution
for the electrolytic Ni deposition. The chemical compositions
of the solutions are reported in Table 1. The procedures
for preparing the solutions were suggested in the literature
[7-10], even if the composition for electroless Ni-B
deposition was optimized after several experiments, because
the deposition was made without a preliminary Zn conversion
coating on the samples, which is frequently used to improve
electroless deposit adhesion.
Diffusion heat treatment of the plated specimens were carried
out in a tubular furnace at 500° C and 530°C for different
times in inert atmosphere. The coating morphology and
alloy microstructure were characterized by optical microscope
and by a Cambridge Stereoscan 440 SEM equipped
with Philips PV9800 EDS. The specimen concentration
profiles were obtained by a LECO GDS 750 A glow dischar-
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TRATTAMENTI TERMICI
7-8/2006 Memorie
Fig. 1 – Ni electrolytic coating on Al
substrate.
Fig. 1 – Rivestimento di Ni elettrolitico sul
substrato di Al.
ge optical spectrometer (GDOS) with an anode surface of 4
mm 2 at 700 V with a current of 20 mA. The microhardness
profile were obtained by a Leitz-Werlag microhardness tester
with 100 g weight. The tribological behavior was investigated
by dry sliding tests with a slider-on-cylinder tribometer.
The stationary sliders consisted of the plated samples
under investigation, whereas the rotating countermaterial
was a hard chromium plated steel cylinder (thickness of the
Cr coating: 300 µm; micro-hardness: 900 HV 0.3
; roughness
R a
: 0.1 µm).
The tests were carried out at 5 N applied load with a sliding
speed of 0,6 m/s and sliding distance of 5000 m at room
temperature. Both friction resistance and system wear (i.e.
wear of the slider plus wear of cylinder) were continuously
measured, using a bending load cell and a LVDT respectively,
and recorded as a function of sliding distance. Values
of the wear scar depths and widths on the slider and the cylinder
were also evaluated at the end of each test by stylus
profilometer. Wear tracks were examined by SEM and OM.
RESULT AND DISCUSSION
Ni electrolytic deposition
Optical microscopy analysis of the electrolytic coating
showed that the layer, constituted by metallic nickel, was
uniform, adherent and thick about 20-40 µm (fig.1).
The diffusion heat treatments, performed at 530° C for 4,
12, 24 h, induced the migration of Ni into the substrate and
the diffusion of the alloying elements towards the surface
with the formation of a diffusion layer about 40-60 µm
depth, as it is shown in the fig. 2.
Increasing the time of heat treatment the thickness of diffusion
zone increased and became less uniform (fig.3). EDS
analysis showed that after 12 h heat treatment, the zone of
diffusion was constituted mainly by Al (82% at.) and Ni
(18% at.). After 24 h heat treatment the composition was not
homogenous and it was possible to distinguish three different
zones (fig.3). The dark thin surface one of about 5 µm
was constituted mainly by Al (81 % at.) and Mg (17.39 %
at.) and only by 0.77 % at. of Ni. In the intermediate layer,
thick about 50 µm, it was detected a constant amount of
20% at. of Ni with a variable composition of Al and Cu. The
concentration of Al and Cu was 64% and 2% in the light
grey phases, while 77% and 14% in the dark grey phases.
The atomic ratio Al/Ni was near to 3, suggesting that Al 3
Ni
was formed. At the interface with substrate there were some
little areas rich in Fe (4% at). The hardness profiles of the
heat treated at 530°C for 12-24 hours were reported in Fig.
4. The sample 12 h treated showed a surface hardness of 770
HV 100
which rapidly decrease to about 300 HV 100
in 50 µm
which correspond to the depth of diffusion layer, suggesting
that the presence of Ni is the responsible of increase of hard-
Fig. 2 – Ni electrolytic coating after heat
treatment at 530°C for 12 h.
Fig. 2 – Rivestimento di Ni elettrolitico dopo
trattamento termico a 530° C per 12h
Fig. 3 – Ni electrolytic coating after heat
treatment at 530°C for 24 h.
Fig. 3 – Rivestimento di Ni elettrolitico
dopo trattamento termico a 530° C per 24h.
Fig. 4 – Hardness profiles of samples heat treated at 530°C for
different times.
Fig. 4 – Profilo di durezza dei campioni trattati a 530° C per
diversi tempi.
Fig. 5 – Ni
electrolytic coating
after heat
treatment at 500°C
for 24 h.
Fig. 5 -
Rivestimento di Ni
elettrolitico dopo
trattamento termico
a 500° C per 24h.
ness. For the samples 24 h heat treated a surface hardness of
180 HV 100
was recorded, which increases to about 800
HV 100
at 50 µm depth. This because, as showed form SEM
analysis, the outer layer is rich in Al and Mg migrated from
the bulk to the surface, while the intermediate zone is constituted
mainly in Al 3
Ni, as revealed by EDS analysis, suggested
that this intermetallic is the main responsible of the
hardening of this layer. Moreover the high spread of the
hardness values was ascribable to the different amount of
Cu founded by EDS analysis and in particular the highest
values have been measured at about 40 µm from the surface
where the concentration of copper is lower. Obviously the
longer diffusion time produces a thicker diffusion layer.
Because both the front of diffusion and profile of hardness
were not uniform it has been chosen to diminish the temperature
of heat treatment from 530°C to 500°C, to limit the
phenomena of element migration from substrate to surface,
maintaining the diffusion time of 24 h.
Fig. 5 showed the sample treated at 500°C for 24 h, where is
present an unique uniform and compact zone of diffusion of
38
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TRATTAMENTI TERMICI
Fig. 6 – Ni
electrolytic coating
of about 80 µm
after heat treatment
at 500°C for 24 h.
Fig. 6 –
Rivestimento di Ni
elettrolitico di 80
µm dopo
trattamento termico
a 500° C per 24h.
Fig. 7 – Hardness profiles of samples heat treated at 500°C for 24 h.
Fig. 7 – Profilo di durezza dei campioni trattati a 500° C per 24h.
Fig. 8 – GDOS depth profiles of NiB coating on Al substrate.
Fig. 8 – Profilo GDOS del rivestimento NiB sul substrato di Al.
about 60 µm which from EDS analysis resulted to be constituted
mainly by Al 3
Ni. The hardness profile of the samples
are reported in fig.7. The values of hardness are almost constant
at 800 HV for about 50 µm due to the presence of
Al 3
Ni.
Samples with a coating of about 80 µm, twice of the others
samples studied, showed, after 24 h heat treatment at 500°
C, two different compact layers of diffusion: one outer zone
with a thickness of about 10 µm and one inner zone with
thickness of about 80 µm (Fig. 6).
From EDS analysis the outer zone resulted constituted
mainly of Al 3
Ni 2
, while the main constituent of the inner zone
was Al 3
Ni. The hardness profiles showed that in proximity
of the surface the hardness values of 1000 HV were
reached, which correspond to the hardness of Al 3
Ni 2
[11].
The value of 800 HV of the inner zone is ascribable to the
presence of Al 3
Ni as main constituent.
Ni-B electroless deposition
Analysis of the electroless coating revealed that the plating
bath allowed the deposition of an amorphous, uniform and
dense Ni-B layer on the 7075 Al alloy.
Fig. 9 – NiB
coating after heat
treatment at 530°C
for 24 h.
Fig. 9 –
Rivestimento di
NiB dopo
trattamento termico
a 530° C per 24h.
The GDOS analysis showed that the layer was about of 10
µm with an average boron content of 1 wt %, (Fig.8).
The heat treatment were carried out at 530°C for different times.
Until 24 h the diffusion was limited. After 24 h heat
treatment an irregular diffusion zone between the coating
and the substrate, thick only about 10 µm, was observed
(fig.9). From EDS analysis, the diffusion zone was constituted
mainly by Al with a ratio Al/Ni near to 3. In the white
spots of fig.9, the presence of a relevant amount of Cu (6%
at.) has been recorded, which was migrated towards the surface
during heat treatment. It was observed that the Ni diffused
along boundary grain. In fact, isles with a different concentration
of Ni (12 % and 6% at., at a depth of 30 µm and
100 µm, respectively) were found until a depth of 100 µm
from the surface, while only trace of Ni was detected around
the isles of diffusion, even if near to the surface layer of diffusion.
An increase in hardness has been recorded in correspondence
of the thin layer of diffusion and of the Ni-rich
isles, where values of about 550 HV 100
and 250 HV 100
have
been measured, respectively. The boron diffusion seems to
not influence the hardness of the substrate, unlike in titanium
alloys [6].
Tribological tests
The first set of tests was carried out at 5N load with a sliding
speed of 0,6 m/s for a sliding distance of 5000 m. Untreated
samples have been compared with Ni coated and heat treated
(at 500° for 24 h) samples. The coefficient of friction of
untreated samples as a function of sliding distance displayed
a wide range of variation with an average value of about 0.5,
while the coefficient friction of all the treated samples
showed a smoother trend and an average value of about 0.7
(fig.10). The values of the wear scar depths on the untreated
samples were about of 300 µm, 30 times greater than those
measured on treated samples, because the surface hardness
of treated samples is higher than untreated ones (fig. 11).
The lower hardness can also account for the lower average
value of the coefficient of friction of untreated samples: the
ploughing component of sliding friction is lower for soft untreated
samples than for hard treated samples. The wide variation
range in the coefficient of friction of untreated samples
can be due to stick/slip motion and extensive grooving
of the soft slider, whereas the smooth trend of the coefficient
of friction of treated samples is due to gradual micropolishing
of the hard slider.
The presence of adhesive yielding and cracks was not observed
in treated samples, because the stronger bond
between coating and substrate after heat treatment, due to
the formation of intermetallics, ensures higher deposition
adherence.
CONCLUSIONS
Heat treatments of Ni electroplated on 7075 Al substrate for
different times and temperatures allowed the interdiffusion
of Ni and element of the constituted principally of Ni and Al
Memorie
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TRATTAMENTI TERMICI
7-8/2006 Memorie
Fig. 10 – Coefficient of friction of Al 7075 T6 and samples Ni
coated and treated at 500°C for 24 h.
Fig. 10 – Coefficiente di frizione della lega Al 7075 T6 e del
campione nichelato elettrolit. e trattato a 500° C per 24 h.
alloy with the formation of a hard diffusion layer. The better
results have been obtained after 24h of heat treatment at
500°C which led to a more uniform zone of diffusion
Surface hardness higher than 1000 HV and depth of hardening
of about 100 µm were achieved by diffusion of Ni layers
and formation of Al 3
Ni 2
and Al 3
Ni intermetallic phases.
For Al alloys the presence of boron not increase the hardness
of the substrate, and electroless coatings are not so efficient
in hardening like electrolytic coatings, because their
thickness is lower.
The tribological test on Al7075 alloy against a hard Cr coated
steel cylinder in air showed an average coefficient of friction
of about 0.5 with a wide variation range, while all the
coated and treated samples exhibited a coefficient friction of
about 0.7. However, the wear scar depth of the hardened alloy
is about 30 times lower than that of the Al7075 alloy.
REFERENCES
[1] S. Tomida, K. Nakata, Surf. Coat. Technol., 174-175,
2003, p. 559
INDURIMENTO SUPERFICIALE DELLA LEGA AL 7075
MEDIANTE TRATTAMENTI DI DIFFUSIONE DI RICOPRIMENTI
DI Ni
Parole chiave: trattamenti termici,
alluminio e leghe, microscopia elettronica, tribologia
In questo lavoro è stato studiato un processo di indurimento
superficiale per la lega Al 7075 attraverso deposizione di
uno strato di Ni e successiva diffusione termica.
I rivestimenti di Ni con spessori variabili tra 10 e 80 _m sono
stati ottenuti mediante un processo di deposizione di Ni
elettrolitico o di Ni-B chimico (tab.1). I trattamenti termici,
il cui scopo è quello di indurre la diffusione del Ni all’interno
del substrato, sono stati effettuati in forno in presenza di
atmosfera inerte alle temperature di 500° e di 530° C per
tempi variabili. L’effetto della temperatura e del tempo di
trattamento termico sull’indurimento superficiale della lega
è stato investigato mediante SEM, EDS, GDOS, prove di
microdurezza e tribologiche.
I trattamenti termici condotti sulla lega di alluminio con lo
strato di Ni elettrolitico ha permesso la interdiffusione del
ABSTRACT
Fig. 11 – Scar depth of Al 7075 T6 and samples Ni coated and
treated at 500°C for 24 h.
Fig. 11 – Profondità del solco della lega Al 7075 T6 e del
campione nichelato elettrolit. e trattato a 500°C per 24 h.
[2] S. Tomida, K. Nakata, S. Saji and T. Kubo, Surf. Coat.
Technol. 142-144, 2001, p. 585
[3] E. Lugscheider, G. Krämer, C. Barimani, H. Zimmermann,
Surf. Coat. Technol, 74-75, 1995, p. 497
[4] K. T. Rie, A. Gebauer, J. Wöhle, Surf. Coat. Technol.,
86-87, 1996, p. 498
[5] G. W. Malaczynski, A. H. Hamdi, A. A. Elmoursi and
X. Qiu, Surf. Coat. Technol., 93, 1997, p. 280
[6] M. Dabalà, K. Brunelli, R. Frattini, M. Magrini, Surf.
Eng., 20, 2004, p. 103
[7] Y. M. Chow, W. M. Lau, Z.S. Karim, Surf. Inter. Anal.,
31, 2001, p. 321
[8] A. R. Giampaolo, J.G. Ordorez, J. M. Guglielmacci, J.
Lira, Surf. Coat. Technol., 89, 1997, p. 127
[9] F. Delaunois, J. P. Petitjan, P. Lienard, M. Jacob-Duliere,
Surf. Coat. Technol., 124, 2000, p. 201
[10] W. D. Fields, R. N. Duncan, J. R. Zickgraf, Metals
Handbook, 9th edn , ASM International, Materials Park
(OH), 1982, p. 601
[11] M. A. Zamzam, A. S. El –Sbagh, M.M. Milad, Materials
and Design, 23, 2002, p. 161
Ni e degli elementi del substrato portando alla formazione
di uno strato esterno indurito costituito principalmente da
Ni e Al. I risultati migliori sono stati ottenuti con il trattamento
termico a 500° C per 24 h che ha permesso la formazione
di uno stato di diffusione compatto ed uniforme costituito
da Al3Ni2 e Al3Ni (fig. 5-6). Con questo trattamento
sono stati ottenuti valori di durezza di circa 1000HV100 per
uno spessore di 100 _m (fig. 7). Temperature maggiori producono
uno strato di diffusione meno uniforme e con valori
di durezza inferiori (fig. 3-4).
Per le leghe di alluminio la presenza di boro non contribuisce
ad aumentare la durezza del substrato e inoltre la deposizione
chimica risulta meno efficace nell’indurimento della
lega rispetto a quella elettrolitica in quanto i rivestimenti
ottenuti presentano uno spessore inferiore (fig. 9).
Le prove tribologiche condotte con un cilindro di acciaio rivestito
al cromo mostrano un valore del coefficiente di frizione
di 0,5 per la lega Al7075 e di 0,7 per i campioni nichelati
trattati termicamente (fig. 10). Tuttavia la profondità dei solchi
dei campioni sottoposti al trattamento di indurimento superficiale
risulta inferiore di circa 30 volte rispetto alla lega
Al 7075 per le prove condotte con un carico di 5 N (fig. 11).
40
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TRATTAMENTI TERMICI
Intelligent control system
for gaseous nitriding process
J. Ratajski, R. Olik, J. Tacikowski, T. Suszko, O. Lupicka
In the present paper, the solutions of the following research issues have been presented:
• dependencies between the process parameters and the layer structure have been defined, which serve to
develop software for the control system of the process, with the view of obtaining a complex layer
structure and optimal kinetics of its creation and growth,
• assumptions for the control system of the gaseous nitriding process have been developed on the basis of
a complementary cooperation of the mathematical model and the indications of the magnetic sensor
registering the nucleation and growth of the layer.
These issues comprise two complementary ways towards the construction of intelligent control systems.
The first one of them consists in developing innovative databases, expert systems to aid the operator in
decision making regarding the choice of defined changes to the process parameters. This is connected
with the knowledge of possibly all the factors and mechanisms which have an influence on the result of
the process and which make it possible to design an algorithm of changes of process parameters that
could guarantee optimal kinetics of the layer growth and its required structure. The other one additionally
makes use of specially constructed sensors, which having been placed directly in the processes react to
the growth of layers and their structure.
Memorie
Key words: mathematical model, model of the nitriding process, system of visualization of nitriding process, control system
INTRODUCTION
A growing range of applications of the nitriding process, in
which the layer nitrided should be characterized by better
and better properties, requires the use of intelligent control
systems. There are two complementary manners of behaviour
towards the construction of such systems. The first
one consists in developing modern databases and expert systems
to support the operator in decision making concerning
the selection of specific changes in the parameters of the
process. It is connected with the knowledge of possibly all
the factors and mechanisms which have an influence upon
the result of the process, which make it possible to design
algorithms of changes of process parameters, which guarantee
an optimal kinetics of the layer growth and its required
structure. The other one additionally uses specially constructed
sensors, which placed directly in the process react
on the growth of layers and their structure.
The issues discussed in this paper include these two manners
of behaviour. In particular, they comprise the following:
1) a development of an experimental-theoretical model of
the nitriding process to facilitate a simulation of the layer
growth kinetics,
2) guidelines for assumptions for the construction of the
following:
• a system for the visualization of growth of the nitrided
layer, whose principle of operation is based on a complementary
interaction of a process result sensor (a magnetic
sensor) and the experimental-theoretical model,
Jerzy Ratajski, Roman Olik, Tomasz Suszko, Oliwia Lupicka
Technical University of Koszalin, Poland
Jan Tacikowski
Institute of Precision Mechanics, Warsaw
Paper presented at the 2 nd International Conference
HEAT TREATMENT AND SURFACE ENGINEERING IN AUTOMOTIVE APPLICATIONS
organised by AIM, Riva del Garda, 20-22 June 2005
• a system of an automatic control of the nitriding process,
in which the block controlling the main process
parameters, i.e. temperature, composition and the intensity
of the flow of the nitriding atmosphere works
on the basis of sensor signal-time courses registered by
a magnetic sensor, mapping the growth of the nitrided
layer.
The model was constructed on the basis of a mathematic description
of the growth kinetics of the nitrided layer and experimental
databases combining the process parameters
with the structure of the nitrided layer. The mathematical description
included the following:
- the growth kinetics of the nitrided layer in quasi-equilibrium
conditions, depending on the process time and nitride
potential,
- the growth kinetics of the diffusion layer.
The assumptions presented in this paper concerning the system
of visualization of the nitrided layer growth and the
control system of the layer growth include innovative solutions
of a new generation. At present, an important issue in
the gas nitriding technology is the possibility to control solely
the process parameters, i.e. the process temperature and
the composition of the gaseous atmosphere and the intensity
of its flow. The creation and the growth of the layer is not directly
controlled in the process. The process, due to the reasons
we know about from practice, may proceed not in compliance
with the assumptions. In such a case, we learn about
the improper course of the layer creation only after its completion.
The assumptions for innovative systems presented
in the paper are based on the indications of the process result
sensor: a magnetic sensor. In particular, the operating principle
of the nitrided layer growth visualization system consists
in a complementary interaction of the system of the sensor
and the model of the process. The characteristic points of
sensor signal-time courses registered by the sensor, which
correspond to e.g. obtaining the maximum surface hardness
or the creation of a continuous layer of iron (carbo)nitrides
(of a few micrometers), constitute the initial conditions for
the description by the model of further stages of the nitrided
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TRATTAMENTI TERMICI
7-8/2006 Memorie
layer development. In the control system and, the software
of the block controlling the main process parameters, i.e.
temperature, the nitriding atmosphere composition and the
intensity of the flow serves to assure that the kinetic sensor
signal-time course registered by a magnetic sensor imaging
the nitrided layer growth is compliant with the reference
course.
MATHEMATICAL MODEL
The course of the kinetics of the process depends on the
speed of particular partial stages. An equation describing the
kinetics of its slowest stage is an equation which describes
the kinetics of the whole process. The speed of particular
stages of the creation and growth of the nitrided layer changes
together with the time of the process. A chemical reaction
of ammonia dissociation on the surface of the nitrided
element is the slowest stage in the initial phase of the layer
creation. While the layer is growing, the diffusion of nitrogen
in steel is becoming the slowest stage. This results in the
fact that the kinetics of the whole process is a complex time
function. However, through a proper selection of the following
parameters: temperature, the intensity of the flow of
the nitriding atmosphere and its composition, one can make
the transport of nitrogen in steel [1] play the main role in the
whole process. Experimental data collected for such cases
concerning the growth of particular mono-phase zones in
the nitrided layer [2], as well as the growth of such zones in
other diffusion systems, e.g. metal-metal [3], indicate a parabolic
low of the growth of phase zones:
where:
∆x i
– thickness of ith phase in n-phase layer after process time
t (for nitrided layer, the maximum value n=3),
k i
– kinetic parameter of the growth of ith phase, the socalled
constant of parabolic growth of ith phase,
t – process time.
Equation (1), with experimentally designated constant k i
,
serves to describe the growth kinetics of phase zones in a given
temperature. For practical reasons, this is enough in most
cases, as the knowledge of k i
value in different temperatures
for all the phases of the diffusing system makes it possible,
in accordance with equation (1), to determine the change
of the thickness of a given ∆x i
phase in the process time
function. However, this is an oversimplified description of
the growth kinetics, which makes it impossible to determine,
in the case of a nitrided layer, e.g. nitrogen diffusion
coefficients in individual layer phases or nitrogen concentrations
on phase borders. One cannot foresee on the basis of
this equation the influence of the remaining phases creating
the layer upon the speed of the growth of a given phase,
either. For this reason, it became necessary to develop general
mathematical equations describing dependencies
between growth parameters (k i
) and diffusion parameters,
which facilitated the determination of diffusion coefficients
in mono-phase layer zones, as well as forecasting of the
phase growth in the function of the process time. In the model
developed, whose particulars were given in paper [4],
the final result are the equations given below (2), which in a
direct manner connect the kinetic parameter of a given phase
k i
with the difference of concentrations on the phase borders
and an effective diffusion coefficient in the phase:
where:
(1)
(2)
k i
– kinetic parameters of the growth of ith and jth phase,
∆c i
– difference of concentrations of the diffusing element
on the border of i th phase.
This is a system of non-linear equations, which can be solved
with e.g. a method of successive approximations (iterations).
In the method of successive approximations, e.g. for zone ε,
the following dependencies are obtained:
From the system of equations obtained it is evident first of
all that the growth parameters are the greater the greater the
difference of concentration of the borders of phases ∆c i
is, in
accordance with the diagram of the phase equilibrium, as
well as the greater the effective diffusion coefficient in a given
phase – (D c i ) ef is.
In the case of alloy ferrite, which includes alloy elements
with a smaller enthalpy of creation of nitrides than iron, an
important objective for mathematical modelling is, apart
from formulae describing a change of the layer thickness,
the determination of the profiles of nitrogen dissolved in the
set of matrix and nitrogen bound in nitrides of alloy elements.
Both these profiles form the distribution of hardnesses
in the nitrided layer. As it is well-known, alloy elements
which form nitrides are characterized by a different chemical
affinity to nitrogen. If there is a strong affinity between
these elements (Ti, V and Cr above 2.5 wt.-%) and nitrogen,
then the front of the reaction of the creation of nitrides is
ahead of the front of nitrogen diffusion [5-8]. Consequently,
a sharp inter-phase border is created between the layer and
the core. In this case, the analysis of the kinetics of the
growth of the diffusion layer amounts to calculations of the
growth of its thickness, while making use of the model developed
for the processes of internal oxidation [9-11]. In this
model, the diffusion of alloy elements is neglected and it is
assumed that there is a constant concentration of nitrogen on
the surface, in equilibrium with the nitriding atmosphere.
With these assumptions, the thickness of the diffusion layer
may be described with the following relationship:
where:
ξ – thickness of diffusion layer,
[N] – surface concentration of nitrogen in at. %,
[X] – concentration of alloy element in at. %,
r – ratio of number of nitrogen atoms to the atoms of alloy
element in nitride,
D – nitrogen diffusion coefficient in ferrite,
t – process time.
In the presence of alloy elements, which are characterised
(3)
(4)
(5)
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TRATTAMENTI TERMICI
7-8/2006 Memorie
cess parameters. Also, a comparison of the distributions of
the total of nitrogen concentrations (Fig.1b), with the experimental
data of hardness (Fig. 1c) serves to indicate an explicit
similarity in their courses for the first and second stages
of the process. This testifies that with the help of the model
developed, one can also simulate in the correct manner
the growth kinetics of the diffusion layer.
THE CONTROL SYSTEM TO THE NITRIDING PROCESS
In the new generation control system to the nitriding process
it was foreseen, independently of the solutions obtained so
far concerning the control of the nitriding atmosphere, a
complementary interaction of the magnetic sensor and the
model of the process. The magnetic sensor constitutes an
element of the system mapping, directly in the process, the
creation and growth of the layer. This system monitors the
nucleation and the growth of the layer nitrided as a result of
a change of the value of the magnetic permeability of the
ferromagnetic diffusion layer, and the creation of iron nitrides
being paramagnetic in the temperatures of the process.
In accordance with the idea of the functioning of the system,
the model of the process plays a double role. As priority, on
the basis of a simulative analysis of the layer growth with
the assumption being made concerning its specific structure,
this model serves to determine the most advantageous algorithm
of the changes of the process parameters. The second
important purpose of the process model for the automatics
system is obtaining information, in the course of the process,
concerning the development of the nitrided layer. In
particular, the model facilitates a current presentation of the
following:
- profiles of nitrogen dissolved in the alloy ferrite and the
profiles of nitrogen bound in nitrides of alloy elements,
- a total profile of nitrogen in the alloy ferrite, mapping qualitatively
the hardness profile, and
- concurrently with the magnetic sensor, it informs about
the thickness of the layer of (carbo)nitrides.
These are very important data which characterise the layer
nitrided, and which condition its usable properties. The model
is the only tool with which one can obtain these. Only
the thickness of the compound layer is concurrently determined
by the system of the magnetic sensor. However, it is a
well know fact that already at the mathematical description,
which constitutes the basis of the experimental-theoretical
model of the process, if the growth kinetics for the nitrided
layer on iron, difficulties are met which are the result of e.g.
lack of possibilities for an accurate and experimental determination
of the values of coefficients of nitrogen transfer
over the inter-phase boundary of gas – metal. The knowledge
of the values of these coefficients is especially important
on the initial stage of the process, in which the reaction of
ammonia dissociation on this boundary constitutes a factor
which determines the speed of the whole process. For this
reason, mathematical descriptions usually do not cover the
initial stage of the layer creation, as it is assumed that its duration
is so short that the description of the kinetics may be
limited to the layer which has already been formed, and
which has constant and equilibrium concentrations on its
boundaries. In the case of iron, neglecting this stage has
practically speaking no impact upon the differences between
the calculated and experimentally determined growth of the
thickness of the diffusion layer, and only introduces a constant
error between the real and calculated thicknesses of the
zones of the nitrided layer, while on steel the conditions for
the process on the stage of the layer creation are most often
decisive for its further correct growth. In the light of the
above, such a solution of the analysis of the layer growth kinetics,
in which the initial stadium is neglected, results in di-
Fig. 2 – Block diagram of the automatics system of the nitriding
process with the visualization system for the course of the layer
growth.
Fig. 2 – Schema a blocchi del sistema automatico del processo di
nitrurazione con il sistema di visualizzazione per il corso della
crescita di strato.
screpancies between the actual course of the layer creation
and the one forecasted in accordance with the model.
In the automatics system discussed, the imperfections of the
model are eliminated by means of a complementary interaction
of the magnetic sensor system, which registers sensor
signal-time courses mapping in a direct manner the growth
process of the nitrided layer. The characteristic points of these
courses correspond with e.g. obtaining the maximum surface
hardness or the creation of a continuous (a few micrometers’)
layer of iron (carbo)nitrides. They constitute the
initial conditions for the description by the model of further
stages of the development of the layer nitrided. Fig. 2 presents
a block diagram of the system, in which the magnetic
sensor together with the process model constitute an independent
module which facilitates a visualization of the course
of the layer growth. This block diagram also includes a
system of an automatic control of the nitriding process, in
which the software of the block controlling the main process
parameters, i.e. the temperature, the composition and intensity
of the flow of the nitriding atmosphere serves to ensure
that the kinetic sensor signal-time course is registered with a
magnetic sensor, mapping the nitrided layer growth, is compliant
with the reference course.
In order to present the visualisation system in a detailed
manner (Fig. 3), a single-stage process of 4340 (AISI) steel
nitriding was selected and conducted in the temperature of
833K and at the nitriding potential K N
= 1.19.
Article [17] presents an analysis of the course of a layer
creation on the example of iron nitriding, which was registered
with a magnetic sensor. The kinetic sensor signal-time
courses obtained in this case should be treated as model
courses, which constitute a reference point for the interpretation
of courses registered during nitriding of steel. Every
steel owing to its specific chemical composition and phase
structure is characterized by an individual susceptibility of
magnetic properties to changes being the result of the creation
of the nitrided layer. Also, processes accompanying the
layer development, such as e.g. the dispersion of the nitrides
of alloy elements being created or their coagulation, which
have an influence, among other things, on the stress state in
the diffusion layer. Consequently, the form of the sensor signal-time
courses depend on the features characteristic of a
given steel. Of course, the mapping proven:
increase /decrease of compressive stresses ⇒ decrease
/ increase of magnetic permeability ⇒ decrease
/ increase of the sensor signal
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TRATTAMENTI TERMICI
Memorie
Fig. 3 – Presentation of the work of the visualizing system for the nitrided layer growth on the example of nitriding 4340 (AISI) steel.
Fig. 3 – Presentazione del lavoro del sistema di visualizzazione dello sviluppo dello strato nitrurato sull'esempio di nitrurazione dell'acciaio
4340 (AISI).
is correct for every ferromagnetic steel, yet in every case it
exerts a different influence on the form of the course registered
by the magnetic sensor.
One of the problems which occur during the control of industrial
nitriding processes with a magnetic sensor is the
so-called temperature background. It is an influence of
changes of the electromagnetic properties of a steel sample
on the registered signal of the sensor, produced by tempe-
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TRATTAMENTI TERMICI
7-8/2006 Memorie
rature changes, which occurs while the furnace is being
heated to the process temperature, and during the execution
of a few-stage process the moment one process stage
goes to another. The temperature background is an individual
feature of every steel and should be subtracted from
the indications of the sensor by the proper software of the
system. As a result, while the furnace is being heated till
the moment the layer starts to create, which initiates the
start of ammonia dissociation (ca. 623K), the signal registered
by the sensor does not change (Fig. 3). Then, atom
nitrogen generated as a result of ammonia dissociation,
diffuses to the interior of the sample and by generating
stresses in it, causes a reduction of the sensor signal. Together
with the registration of the sensor signal-time courses,
the software of the sensor system facilitates a presentation
of its differential signal, whose extreme points indicate
the occurrence of characteristic layer creation stages,
like e.g. the beginning of the creation of surface iron nitride
phases (Fig. 3). This is documented by an X-ray analysis
of nitrided samples together with the sample located in
the sensor and samples individually taken out from the retord,
in characteristic moments of the process evident from
the indications of the sensor [17]. On the basis of the differential
course, one can also obtain information on the phase
composition and the nitride layer created in the process
and located by the surface. As it is know, [18,19] ε and γ’
phases, which are paramagnetic in the temperatures of the
process, are distinguished by a different Curie temperature,
i.e. the magnetic transformation point - γ’ phase goes into a
ferromagnetic state in the temperature of 763K, and ε
phase in the range of temperatures of 303-593K. For this
reason, the course of the sensor signal, registered during
the cooling down stage, has two maximum values being
the result of magnetic transformations of the abovementioned
phases. The area limited by these maximum values is
proportional to the volume fraction of ε and γ’ phases in
the layer of nitrides.
Together with the indications of the sensor, which map the
growth of the nitrided layer, calculations are conducted with
the help of the model of the process, which represent the development
of profiles of nitrogen dissolved in the alloy ferrite,
and of nitrogen bound in the nitrides of alloy elements,
the degree to which the surface is covered with iron (carbo)nitrides,
and in the further part of the process, informing
on the changes of the thickness of the (carbo)nitride layer
depending of the process time.
SUMMARY
The model of the nitriding process presented in this paper
will serve to determine the procedures of a comprehensive
simulation of the process. It covers the growth kinetics of
mono-phase zones in the nitrided layer on iron and low-carbon
steels, as well as the calculation of nitrogen profiles in
the diffusion layer on alloy steels. In particular, the model
proposed, describing the growth of individual zones of the
nitrided layer in the conditions of a quasi-equilibrium, is a
simple and convenient tool facilitating both the determination
of the nitrogen diffusion coefficients in mono-phase
layer zones, as well as forecasting their growth in the function
of the process time and the nitrogen potential. In the
case of a steel including alloy elements with a smaller
enthalpy of the generation of nitrides than iron, the model
facilitates a calculation of nitrogen concentrations, which
constitute a starting point for the development of procedures
enabling forecasting of hardness profiles on the basis of the
process parameters assumed and the chemical composition
of steel.
The system of an automatic control of the process presented
in the paper constitutes an innovative and unique solution.
In the new control system, independently of the so-far existing
solutions concerning the control of the nitriding atmosphere,
interactions of the magnetic sensor system and the
process model have been foreseen. The magnetic sensor is a
component of the system which maps directly in the process
the creation and growth of the layer. In turn, the model serves
to determine the most advantageous algorithm of the
changes of the process parameters, and to provide complementary
to the indications of the magnetic sensor, information
concerning the growth of the nitrided layer.
The issues analysed in this paper and the results obtained
constitute elements of the construction of modern and effective
control systems, in which artificial intelligence elements
may be applied. These systems make it possible to face
new challenges connected with the applications of the nitriding
process.
REFERENCES
1) I. TORCHANE, P. BILGER, J. DULCY, M. GANTOIS,
Control of Iron Nitride Layers Growth Kinetics in the
Binary Fe-N System, Metall. Trans., A 27 (1996),
p.1823-1835,
2) W. PITSCH, E. HOUDREMONT, Ein Beitrug zum System
Eisen-Stickstoff, Archiv für das Eisenhüttenwesen,
27 (1956), p.281-284.
3) K.P. GUROW, B.A. KARTASZKIN, J.E. UGASTIE,
Wzimnaja Diffuzija w Mnogofaznych Sistemach, Moskwa
"Nauka", (1981).
4) J. RATAJSKI, Model of growth kinetics of nitrided
layer in the binary Fe-N system. Zeitschrift fur Metallkunde,
95 (2004), p.9-23.
5) B. MORTIMER, P. GRIVESON, K.H. JACK, Precipitation
of Nitrides in Ferritic Iron Alloys containing chromium,
Scan, J. Metallurgy, 1 (1972), p.203-209.
6) D.H. JACK, P.C. LIDSTER, P. GRIEVESON, K.H.
JACK, Kinetics of Nitriding Iron Alloys, Scan, J. Metallurgy,
1 (1971), p.374-379.
7) Y. SUN, T. BELL, Mathematical modelling of the plasma
nitriding, Proc. of the 9th Int. Cong. on Heat Treatment
and Surface Engineering, Nice, France, (September
1994), IFHT, p.385-390.
8) Y. SUN, T. BELL, Modelling of plasma nitriding of low
alloys steels, Surf. Eng., 11 (1995), p.146-148.
9) C. WAGNER, Reaktionstypen bei der Oxydation von
Legierungen, Z. für Elektrochemie, 63 (1959), p.772-
782.
10) R.A. RAPP, H,D. COLSON, The Kinetics of Simultaneous
Internal Oxidation and External Scale Formation
for Binary Alloys, Metall. Trans. A 236 (1966), p.1616-
1618.
11) F.N. RHINES, Gas-Metal Diffusion and Internal Oxidation,
Atom Movements, Amer. Soc. Met. Cleveland,
(1951).
12) B.J. LIGHTFOOT, D.H. JACK, H. DU, Kinetics of nitriding
with and without white-layer formation, Proc. of
the Conf. on Heat Treatment ’73, London, (1973), p.59-
65.
13) G. ROBERTS, Diffusion with Chemical Reaction, Metal
Science, 2 (1979), p.94-96.
14) J. CRANCK, The Mathematics of Diffusion, Oxford, at
the Clarendon Press, (1956).
15) J. RATAJSKI, J. TACIKOWSKI, M.A. SOMERS, Development
of compound layer of iron (carbo)nitrides
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during nitriding of steel, Surface Engineering Vol. 19
No. 4, (2003), p.285.
16) J. RATAJSKI, J. TACIKOWSKI, T. SUSZKO, O. LU-
PICKA, Modelling of structure and properties of materials
in the nitriding process, Inzynieria Powierzchni,
No 1, (2005), p.62 (in Polish).
17) J. RATAJSKI, Monitoring nitride layer growth using
magnetic sensor, Surface Engineering Vol. 17 No. 3,
(2001), p.193.
18) K.H. EICKEL, W. PITSCH, Magnetic property of heksagonal
nitride Fe2,3N, Phys.Stat.Sol., 39, (1970),
p.121.
19) B.C. FRAZER, Magnetic structure of Fe4N, Phys. Review,
112, (1958), p.751.
Memorie
ABSTRACT
CONTROLLO DI SISTEMA INTELLIGENTE
PER IL PROCESSO DI NITRURAZIONE A GAS
Parole chiave:
modellazione, processi, trattamenti termici, controlli
Nel presente lavoro vengono presentate le soluzioni dei seguenti
oggetti di ricerca:
- sono state definite dipendenze fra i parametri di processo
e struttura degli strati, al lo scopo di sviluppare il software
per il sistema di controllo di processo, con l’intento di
ottenere una struttura con strato complessa e una cinetica
ottimale di creazione e crescita,
- sono stati sviluppati i presupposti per il sistema di controllo
di processo della nitrurazione a gas sulla base di
una cooperazione complementare del modello matematico
e delle indicazioni del sensore magnetico che registra la
nucleazione e la crescita dello strato.
Questi aspetti contengono due direzioni complementari
verso la costruzione di sistemi di controllo intelligenti. Il
primo consiste nello sviluppo di database innovativi, sistemi
specializzati per aiutare l'operatore nelle decisioni frelative
alla scelta di cambiamenti definiti dei parametri di
processo. Ciò è collegato possibilmente alla conoscenza di
tutti i fattori e meccanismi che influiscono sui risultati del
processo e che permettono di progettare un algoritmo di
cambiamenti dei parametri di processo che potrebbero garantire
la cinetica ottimale della crescita dello strato e della
relativa struttura necessaria. L’altro utilizza sensori
specialmente costruiti, posti direttamente all’interno dei
processi che reagiscono alla crescita degli strati e della loro
struttura.
7-8/2006
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RIVESTIMENTI
Production of aluminium
coated ferritic stainless steel
by co-rolling and annealing
D. Pilone, F. Felli, U. Bernabai
Ferritic stainless steel enriched with aluminium (>7%) represents the best material for the production
of substrates in catalytic converters because of its good corrosion resistance at high temperature
(700-1000 °C). Enriching steel surface with aluminium avoids brittleness problems related
to the traditional metallurgical process and, due to the high superficial aluminium concentration,
enhances the Al 2
O 3
scale formation in the early stages of the oxidation process. Among several
techniques, the co-rolling process appears to be a promising technology for enriching ferritic stainless
steel surface with aluminium. That versatile technology produces, at room temperature,
a very adherent aluminium coating without affecting the structural and mechanical properties
of the substrate. In this work annealed AISI 430 ferritic stainless steel has been cold rolled together with
aluminium foil to form an Al-steel-Al sandwich. The aim of the work was to enhance the oxidation
resistance through a controlled oxidation of a thin aluminium layer. The diffusion bonding obtained via an
annealing treatment was evaluated studying the concentration profiles as a function of process
temperature. Intermetallics and /or solid solutions formed were characterized by X-Ray Diffraction (XRD).
Memorie
Key words: Alumina coating, co-rolling, iron aluminides, solid state diffusion, catalytic converter
INTRODUCTION
D. Pilone, F. Felli, U. Bernabai
Dip. ICMMPM – Università degli Studi di Roma “La Sapienza”, Roma, Italy.
Paper presented at the 2 nd International Conference
HEAT TREATMENT AND SURFACE ENGINEERING IN AUTOMOTIVE APPLICATIONS,
organised by AIM, Riva del Garda, 20-22 June 2005
Ferritic stainless steel enriched with aluminium (>7%) represents
a promising material for the production of substrates
in catalytic converters because of its good corrosion resistance
at high temperature (700-1000 °C) and because of the
absence of nickel, which is responsible for localized corrosion
in sulfidizing environments.
Iron aluminides have been of great interest for many years
because of their excellent oxidation and sulfidation resistance,
as well as because of their reduced density. Despite that,
their use as a structural material has been limited by limited
room temperature ductility and a drop in strength at temperature
higher than 600 °C [1]. Formation of iron aluminides
includes ingot solidification followed by thermomechanical
treatment, but over the past 15 years many attempts have
been done to form iron aluminides by solid state methods
such as hot pressing [2], mechanical alloying [3] or in-situ
synthesis for surface coating on conventional metallic materials
[4]. One promising method is the formation of aluminides
by heat treatment of co-rolled foils [5]. This method has
been used over the past few years to produce Al-Ni intermetallics
by cold rolling and annealing of Al/Ni alternate foils
[6,7]: a reaction in the solid state is promoted by milling together
alternate foils of the elements and by low temperature
annealing (T

RIVESTIMENTI
7-8/2006 Memorie
degreased using ultrasonic cleaning. Steel foil was also mechanically
brushed in order to activate the metallic surface.
Al and steel sheets, having a surface area equal to 25x90
mm, were superimposed in order to obtain an Al-steel-Al
sandwich 80 µm thick. The assemblies were cold-rolled in
air to reach a final thickness of 50 µm. After co-rolling the
assemblies were thermally treated: 1) at 600°C for 12 hours,
2) at 600°C for 12 hours and at 900°C for 40 minutes, 3) at
900°C for 40 minutes.
Polished cross sections of specimens both after reaching the
final thickness and after heat treatment were imaged by
SEM, while EDS analyses were carried out in order to obtain
Al concentration profiles close to Al-steel interface.
Small parts of the assembly were sampled and analysed,
using X-ray diffraction (XRD), to identify different phases
formed as a consequence of heat treatments.
Thermal cycling tests were performed on specimens pretreated
at 600 °C for 12 hours. The chosen thermal cycle
was: 15 minutes either at 700°C or 800°C or 900°C, 5 minutes
at room temperature.
SEM imaging was used to evaluate oxide compactness and
morphology after heat treatment as well as after thermal cycling.
RESULTS AND DISCUSSIONS
Al-steel-Al assemblies co-rolling produced, after few rolling
passes, a good adherence among layers: initial thickness of
different layers and number of rolling passes were determined
with great care in order to avoid specimen embrittlement
due to cold working as well as cracks formation.
After reaching the desired final thickness a SEM analysis of
the specimen cross section revealed that each Al layer was
about 10 µm thick. Since the rolling process produced only
a cold welding of different layers, heat treatments were
performed with the objective of promoting both Al, Fe and
Cr diffusion and growth of an oxide protective layer.
1. Isothermal tests
On the basis of previous results [8,9] the sample was held
isothermally at 600 °C for 12 hours to activate diffusion and
solid state reactions. XRD patterns revealed the presence in
the annealed sample of Al-rich intermetallic compounds such
as FeAl 2
and FeAl 5
(Fig. 1a) together with metallic Al
and alumina. SEM micrograph of the specimen cross section
highlighted that the alumina layer, which appears compact
and well adherent to the substrate, has a thickness of
about 5 µm (Fig. 2 ). If a specimen pre-treated in that way is
kept for 40 minutes at 900 °C, the alloy phases revealed by
XRD are completely different (Figure 1b) from the previous
ones. High temperature treatment, favouring both Al diffusion
toward the specimen centre and Al→Al 2
O 3
conversion,
decreases Al concentration in the outer zone. As suggested
by Al-Fe phase diagram, by lowering Al concentration in the
alloy, α2 phase formation is favoured although a small
quantity of Fe 2
Al 5
is still present together with aluminium
oxide. If the co-rolled assembly is treated at 900 °C without
a preoxidation stage, the phases detected by means of XRD
are alumina, Cr 3
O and a ternary Al-Fe-Cr intermetallic compound
(Fig. 1c). In fact at high temperature chromium easily
diffuses through the scale and reacts with oxygen, moreover
when Cr concentration increases in the outer zone, it forms
intermetallic compounds with Al and Fe.
Treatment temperatures have to be carefully chosen so that
aluminium concentration in the ferritic steel reaches 10-14%
by weight avoiding the formation of intermetallic compounds,
which would determine material embrittlement.
Considering that the treatment temperature affects not only
the base material structure, but also the oxide growth, deter-
Fig. 1 – X-Ray diffractograms of samples after co-rolling and heat
treatment. Key: (a) 12 h at 600 °C, (b) 12 h at 600 °C + 40 min at
900 °C, (c) 40 min at 900 °C. ■ Al 5
Fe 2
orth., ◆ FeAl2, ✱ Al,
▲ Al 2
O 3
,, ✼ FeAl, ● Fe 2
Al 5
mon., ✝ Cr3O,
∞ Al 0.99
Cr 0.02
Fe 0.99
.
Fig. 1 – Analisi XRD di campioni sottoposti a colaminazione e
trattamento termico. (a) 12 ore a 600 °C, (b) 12 ore a 600 °C + 40
minuti a 900 °C, (c) 40 minuti at 900 °C. ■ Al 5
Fe 2
orth.,
◆ FeAl2, ✱ Al, ▲ Al 2
O 3
,, ✼ FeAl, ● Fe 2
Al 5
mon., ✝ Cr3O,
∞ Al 0.99
Cr 0.02
Fe 0.99
.
mining treatment parameters is the key issue in the process
set up. SEM imaging was used to study oxide layer morphology.
As it can be observed in Figure 3a, after keeping the
specimen at 600 °C for 12 h, the oxide appears to be porous
and characterized by ripples and globular morphology. If
that specimen is treated at 900 for 40 minutes the nuclei previously
formed grow with consequent porosity increase
(Fig. 3b). If the cold rolled specimen is treated at 900 °C for
40 minutes, without a pre-oxidation treatment, the oxide
morphology is completely different: a naked eye observation
reveals the presence of a thin and glassy alumina layer,
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RIVESTIMENTI
Memorie
Fig. 2 – SEM micrograph of
the specimen cross section
after heat treatment at 600 °C
for 12 hours,
Fig. 2 – Micrografia SEM
della sezione del co-laminato
dopo trattamento termico a
600 °C per 12 ore.
Fig. 3 - SEM micrographs showing the samples surface morphologies after co-rolling and different heat
treatments: (a) 12 h at 600°C, (b) 12 h at 600°C + 40 min at 900°C, (c) 40 min at 900°C.
Fig. 3 – Micrografie SEM in cui è mostrata la morfologia superficiale dei campioni dopo co-laminazione e
differenti trattamenti termici: (a) 12 ore a 600°C, (b) 12 ore a 600°C + 40 minuti a 900°C, (c) 40 minuti a
900°C.
Fig. 5 – SEM micrographs
showing the samples surface
morphologies after co-rolling,
annealing at 600°C for 12
hours and thermal cycling at:
(a) 700 °C, (b) 800 °C, (c)
900°C.
Fig. 5 – Micrografie SEM in
cui è mostrata la morfologia
superficiale dei campioni
dopo colaminazione, ricottura
a 600°C per 12 ore e ciclaggio
termico a: (a) 700 °C, (b) 800
°C, (c) 900°C.
rous structure is essential to produce both a higher catalyst
effectiveness and a greater oxide adherence due to reduced
internal compressive stresses.
2. Thermal cycling tests
During their life, catalytic converters are subjected to cyclic
thermal variations, which are critical since they can produce
scale spallation reducing the converter life. The specimens
previously isothermally treated at 600 °C for 12 hours, were
subjected to thermal cycling tests at 700 °C, 800 °C and 900
°C. As expected the tendency to scale formation increases
with increasing testing temperature.
Fig. 4 – X-Ray diffractograms of samples after co-rolling, heat
treatment at 600°C for 12 h and 1000 thermal cycles at 700 °C.
▲ Al 2
O 3
, ✻ FeAl, + Cr 3
O, ∞ Al 0.99
Cr 0.02
Fe 0.99
.
Fig. 4 – Analisi XRD di un campione co-laminato dopo
trattamento termico a 600 °C per 12 ore e dopo 1000 cicli in un
test di ciclaggio termico a 700 °C.
▲ Al 2
O 3
, ✻ FeAl, + Cr 3
O, ∞ Al 0.99
Cr 0.02
Fe 0.99
.
which makes the surface really shiny. Figure 3c shows a
SEM micrograph of the specimen surface highlighting the
formation of a compact oxide layer with evident cracks due
to internal stresses arising from metal-oxide transformation.
That oxide morphology is not desirable since a tiny and po-
2.1 Thermal cycling at 700 °C
XRD analysis of the sample after 1000 thermal cycles at 700
°C shows that, as interdiffusion proceeds with time, Al-rich
intermetallics disappear, while α2 phase is formed (Fig. 4).
As far as oxide morphology is concerned, thermal cycling at
700 °C modifies the oxide layer obtained during pre-oxidation
forming a more compact scale. That scale has lower internal
compressive stresses and then it is more susceptible of
deep cracking (Fig. 5a ). Chromium oxide found after thermal
cycling was probably formed in those areas where the
alumina layer cracked, allowing the reaction between
chromium and oxygen.
2.2. Thermal cycling at 800 °C
Thermal cycling at 800 °C determines the growth, on the
oxide formed during the pre-oxidation stage, of a high num-
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RIVESTIMENTI
7-8/2006 Memorie
ber of alumina nuclei that appear acicular in shape (Fig. 5b
). XRD analyses show that the phases in the scale are the same
that were found in the sample after thermal cycling at
700 °C. In this case, because of acicular growth, the compressive
stresses are still present in the oxide layer, so favouring
the resistance to scale cracking.
2.3. Thermal cycling at 900 °C
Thermal cycling at 900 °C produces a rapid oxide growth
and the formed scale appears quite compact. As it can be observed
in Figure 5c the scale, which probably still has internal
compressive stresses, cracks only on the top of the ripples
because of oxide lateral growth.
XRD patterns revealed that thermal cycling at 900 °C determines
the formation, in addition to alumina, of a considerable
amount of chromium oxides, more rich in oxygen in
comparison with the one found after thermal cycling at 800
°C. As a consequence of high temperature treatment, which
promotes aluminium diffusion toward the specimen centre as
well as aluminium oxidation, a Cr-Fe intermetallic is formed.
CONCLUSIONS
Annealed AISI 430 ferritic stainless steel has been cold rolled
together with aluminium foil to form an Al-steel-Al
sandwich in order to obtain a mechanical bonding among
the three layers. After co-rolling the assembly was thermally
treated with the dual objective of producing diffusion bonding
between Al and steel and of obtaining the growth of a
porous and adherent oxide layer. Isothermal tests highlighted
that alumina films formed at 600 °C have a considerable
degree of porosity, while alumina films formed at 900 °C
appear thin and glassy. On the basis of those results a preoxidation
stage at 600 °C appears essential to obtain a porous
oxide coating allowing for relaxation of thermal and
growth stresses. Under thermal cycling conditions the oxide
layer cracks at 700 °C, while it performs much better at 800
COLAMINATI IN FOGLIO SOTTILE
ACCIAIO INOX FERRITICO-ALLUMINIO
Parole chiave: acciaio inox, ossidazione, laminazione,
diffrattometria, microscopia elettrica
Un acciaio inossidabile ferritico contenente più del 7% di
alluminio rappresenta il materiale più idoneo per la produzione
di substrati per marmitte catalitiche in virtù della buona
resistenza alla corrosione alle alte temperature (700-
1000°C). Un arricchimento superficiale con alluminio della
superficie dell’acciaio tipo AISI 430 permette di evitare i
problemi di infragilimento correlati con il processo metallurgico
tradizionale e favorisce la formazione dello strato
superficiale di Al 2
O 3
nei primi stadi del processo di ossidazione.
Tra le varie possibili tecniche di arricchimento superficiale
dell’acciaio con alluminio la co-laminazione appare
la più promettente.
Nel presente lavoro colaminati in foglio sottile sono stati
ABSTRACT
°C since that temperature treatment produces the growth of
a tiny acicular structure, which avoids coating spallation.
Thermal cycling at 900 °C produces a scale consisting of
mixed oxides containing mostly Al and Cr: its morphology,
characterised by ripples, is beneficial for thermal stresses
compensation and assures good adherence between coating
and substrate even at high temperature.
ACKNOWLEDGEMENTS
The authors thank CSM for the financial support.
REFERENCES
[1] C.G. MCKAMEY, J.H. DE VAN et al., J. Mater. Res. 6,
(1991), p.1779.
[2] B.H. RABIN and R.N. WRIGHT, Metallurgical transaction
A 22, (1991), p.277.
[3] G.H. FAIR and J.V. WOOD, Journal of Materials
Science 29, p.1935.
[4] J. DUSZCZYK, J. ZHOU et al., Journal of Materials
Science 34, (1999), p. 3937.
[5] J.R. BLACKFORD, R.A. BUCKLEY et al., Scripta
Materialia 34, (1996), p. 721.
[6] L.BATTEZZATI, P. PAPPALEPORE et al., Acta Materialia
47, (1999), p. 1901.
[7] L. BATTEZZATI, C. ANTONIONE et al. 3, (1995), p.
67.
[8] G.A. CAPUANO, A. DANG et al., Oxidation of Metals
39, (1993), 263.
[9] U. BERNABAI, A. BROTZU et al., Proc. 4th ASM
Heat Treatment and Surface Engineering Conference in
Europe, Florence (1998), AIM, p. 383.
[10] U. BERNABAI and F. FELLI, Proc. Materiali: Ricerca
e Prospettive tecnologiche alle soglie del 2000, Milano
(1997), FAST, p. 1129.
prodotti a partire da fogli di acciaio inossidabile ferritico,
spessi 50 µm, e da fogli di alluminio raffinal aventi uno
spessore di 15 µm. Tali fogli sono stati colaminati allo scopo
di ottenere un sandwich Al-acciaio-Al avente uno spessore
finale pari a 50 µm. Il colaminato è stato trattato isotermicamente
1) a 600 °C per 12 ore; 2) a 600 °C per 12 ore e
a 900 °C per 40 minuti; 3) a 900 °C per 40 minuti. Analisi
XRD hanno rivelato (Fig.1a) la formazione, per effetto della
diffusione e di reazioni allo stato solido, di intermetallici
ricchi in alluminio quali FeAl 2
e Fe 2
Al 5
dopo il solo trattamento
a 600 °C. Il trattamento a 900 °C del campione già
trattato a 600 °C determina non solo la conversione Al →
Al 2
O 3
ma anche la diffusione dell’alluminio verso il centro
del campione: come suggerito dal diagramma di stato Al-
Fe, al diminuire della concentrazione dell’alluminio nella
lega, si favorisce la formazione della fase α2 (Fig.1b). Se il
colaminato è trattato esclusivamente a 900 °C le analisi
XRD rivelano la presenza di ossido di cromo e di un composto
intermetallico ternario Al-Fe-Cr (Fig.1c).
58
la metallurgia italiana

RIVESTIMENTI
Le temperature di trattamento devono essere scelte attentamente
in modo che la concentrazione dell’alluminio nell’acciaio
raggiunga il 10-14% in peso evitando la formazione di
intermetallici fragili. Poiché la temperatura di trattamento
influenza non solo la struttura della lega, ma anche la
morfologia dello strato di ossido, che dovrà fungere da supporto
per il catalizzatore, la determinazione dei parametri
di processo è il punto chiave nella messa a punto del trattamento.
Dopo il trattamento a 600 °C per 12 ore l’ossido appare
poroso e caratterizzato da una crescita globulare
(Fig.3a); se il provino pre-ossidato è trattato a 900°C per
40 minuti i nuclei precedentemente formati si accrescono e
la porosità dell’ossido aumenta (Fig.3b). Se il colaminato è
trattato a 900 °C, senza pre-ossidazione, si forma uno strato
di ossido sottile e vetroso caratterizzato da profonde cricche
(Fig.3c) causate da sollecitazioni interne conseguenti alla
trasformazione metallo-ossido.
Poiché durante l’esercizio i convertitori catalitici sono soggetti
a variazioni cicliche di temperatura, che possono produrre
un distacco della scaglia di ossido, campioni colaminati,
pre-ossidati a 600 °C, sono stati sottoposti a prove di
ciclaggio termico a 700, 800 e 900 °C. Dopo 1000 cicli a
700 °C la scaglia di ossido appare più compatta e fratturata
rispetto a quella ottenuta dopo la pre-ossidazione (Fig. 5a),
mentre dopo 1000 cicli a 800 °C la scaglia di ossido è caratterizzata
da un numero molto elevato di nuclei accresciutisi
con morfologia aciculare (Fig.5b).
Prove di ciclaggio termico a 900 °C hanno determinato la
formazione di una scaglia costituita prevalentemente da ossidi
di alluminio e cromo: la morfologia di tale scaglia, caratterizzata
da protuberanze (Fig. 5c), appare favorevole
per la compensazione degli stress termici ed assicura una
buona aderenza tra substrato e rivestimento anche alle alte
temperature.
Memorie
7-8/2006
la metallurgia italiana 59