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EFFECT OF THE HEAT TREATMENT ON
THE MECHANICAL PROPERTIES OF A
PRECIPITATION HARDENING STEEL
FOR LARGE PLASTIC MOLDS
D. Firrao, P. Matteis, G. M. M. Mortarino, P. Russo Spena, M. G. Ienco, G. Pellati,
M. R. Pinasco, R. Gerosa, G. Silva, B. Rivolta, M. E. Tata, R. Montanari
Continuously growing activity in the area of the engineering plastics led to the necessity of developing
new low-cost, high-performance plastic mold steels. In fact, when it is necessary to fabricate large plastic
components, such as bumpers and dashboards for motor vehicles, the traditionally adopted ISO 1.2738 plastic
mold steel exhibits low fracture toughness and highly inhomogeneous microstructures (continuously varying
from surface to core), as obtained from the pre-hardening (quenching and tempering) of large blooms. New
alloys and alternative manufacturing routes may allow to obtain plastic injection molds with good mechanical,
wear and weldability properties. Precipitation hardening tool steels are being proposed for such an application,
yielding improved mechanical properties and lower overall costs and lead-time. A precipitation hardenable
steel, developed for injection molding of large engineering polymer components, was investigated.
The microstructures and the mechanical properties of the precipitation hardenable steel bloom were
investigated after the steelwork heat treatment. Moreover, the strengthening mechanism by means of aging
heat treatments was examined on samples subjected either to the steelwork heat treatment only, or also to a
successive laboratory heat treatment. To the purpose, X-rays diffraction and EDS analyses were carried out in
order to indentify second phases electrochemically extracted from aged and not aged samples.
KEYWORDS: plastic mold steel, precipitation hardening, metallography, mechanical properties, fracture toughness,
fractography
INTRODUCTION
Large steel molds are employed in injection molding processes to
fabricate massive plastic automotive components (such as bumpers and dashboards), by using glass-reinforced thermoplastic polymers. During the service, several stresses act on a plastic mold:
polymer’s injection pressure, mechanical and thermal fatigue (a
few millions of pieces can be fabricated with one mold), and wear
from reinforced resin flows; stresses can be further enhanced by
notch effects and by abnormal shop operations.
The molds are commonly machined from large quenched and
tempered blooms, typically with 1x1 m cross-section and more
than 1 m length. The ISO 1.2738 (or 40CrMnNiMo8-6-4 [1]) alloy
Donato Firrao, Paolo Matteis, Giovanni M. M. Mortarino,
Pasquale Russo Spena
Politecnico di Torino, Italy
Maria G. Ienco, Gabriella Pellati, Maria R. Pinasco
Università di Genova, Italy
Riccardo Gerosa, Giuseppe Silva, Barbara Rivolta
Politecnico di Milano, Italy
Maria E. Tata, Roberto Montanari
Università di Roma Tor Vergata, Italy
steel grade is the most used steel. Due to the large section, blooms
of the above steel exhibit after, heat treatment, inhomogeneous
microstructures and mechanical properties continuously varying
from the surface to the core of the bloom; impact notch strength
and fracture toughness are everywhere quite low (at the 10 J and
40 MPa√m level, respectively [2]). Moreover, the ISO 1.2738 steel
is difficult to weld (1.16 carbon equivalent index [3]), although
weld bed deposition operations are usually necessary to modify
the mold face, also to extend the service life during model revamping.
Several precipitation hardening steels have been proposed as an
alternative, with the aim of yielding more uniform microstructures and better properties throughout the mold sections, and to
improve weldability (a carbon content lower than 0.4% may be
adopted).
The P21 [4,5] standard grade steel, for example, contains 0.2% C,
4% Ni, 1.2% Co, and lower amounts of V, Al, Mn, Si, Cr [4]; yet,
most grades are proprietary and not disclosed in detail [6]. The
solubilization temperature can be subcritical, as for the P21 grade
[4,7] (albeit after an hypercritical annealing [7]), or hypercritical,
for some proprietary grades, whereas the aging temperature is always subcritical (e.g. 530 °C for the P21 grade [4,7]), and therefore
yields only very limited dimensional variations. The final (servla metallurgia italiana >> aprile 2009
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a
b
s
Fig. 1
Heat treatment schedule (dashed lines represent uncontrolled air cooling stages). a): Preliminary heat treatment (grey
lines, including de-hydrogenation), final heat treatment (black lines). b): Bloom cooling curves from surface to core (constantparameters analytical calculation) superimposed to the CCT diagram.
Trattamento termico complessivo (le linee tratteggiate rappresentano raffreddamenti liberi in aria). a): Trattamento termico preliminare
(linee grigie, includente la de-idrogenazione), trattamento termico finale (linee nere). b): Curve di raffreddamento del blumo da superficie
a cuore (calcolo analitico a parametri costanti) sovrapposte al diagramma CCT.
ice) hardness is usually in the 37-42 HRC [6] range.
The precipitation hardening heat treatment can be performed after mould machining, owing to the fact that it induces only very
limited deformations [8], and a suitable aging treatment may yield
homogeneous microstructures and mechanical properties in geometrically complex and large moulds. Furthermore, in some cases
(e.g. for the here proposed steel), the preliminary heat-treatment
performed on the as-forged bloom can employ air cooling after
austenitization, as opposed to oil quenching in the case of traditional hardened steels, thus yielding much lower temperature
gradients and minimizing residual stresses.
In the present work a precipitation hardenable steel is investigated; its chemical composition is listed in Tab. 1. The aging of
this steel in the 550 to 630 °C temperature range was previously
studied by means of hardness tests performed after increasing aging durations [9], evidencing overaging peaks after 1 or 2 h aging
at 630 or 590 °C, respectively, as opposed to almost asymptotic
aging at 550 °C up to 20 h duration. For this reason, and since
the homogeneous heating of large moulds may require several
hours and implies different actual durations at temperature from
surface to core, only aging temperatures equal or lower than 550
°C are considered here.
EXPERIMENTAL PROCEDURES
The bloom production cycle, performed in the steelwork, consistC
0.05 - 0.15
Mn
0.1 - 1.1
Cr
0.1 - 0.9
ed of several steps: ingot casting, hot forging (in order to reduce
the microstructural and chemical inhomogeneities and to obtain
a 2400 (L) x 1500 (T) x 500 (S) mm bloom), preliminary heat treatment (including a dehydrogenization process), austenitization/
solubilization, air quenching, low-temperature tempering. The
heat treatment schedule is displayed in Fig.1a.
By superimposing the austenite cooling stages (of the bloom heat
treatment) to the steel’s CCT diagram (Fig. 1b), it can be hypothesized that: i) during the first austenitization, most of the primary
Mo and V carbides were dissolved; ii) during the first air cooling
and furnace cooling stages, the austenite-to-pearlite transformation was avoided, a fine and homogeneous carbides re-precipitation occurred, and the austenite was finally transformed into
bainite at lower temperatures, no transformation taking place in
the final uncontrolled air cooling to room temperature; iii) the
de-hydrogenization treatment, performed in two steps, caused
(tempering and) aging of the bainitic matrix and carbides coarsening. Moreover, as it regards the final heat treatment, the bloom
was austenitized at a lower temperature (1020 °C), so that Mo
and V carbides were partially dissolved, and it is hypothesized
that, after the air quenching, a homogeneous and fully bainitic
microstructure was obtained both in the surface and in the core of
the bloom. Finally, the steel was subjected to a double tempering
at 400 °C.
All the examined samples were cut from the bloom at a distance
of less than 170 mm from the bloom surface in the S direction.
Ni
2.5 - 4.5
Mo
2.5 - 4.5
Si
0.1 - 1.1
V
0.05 – 0.20
s
Tab. 1
Nominal chemical composition of the proposed steel (wt. %) (actual analysis covered by industrial confidentiality).
Composizione chimica nominale dell’acciaio proposto (% in peso) (analisi esatta coperta da riservatezza industriale).
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a
b
c
d
s
Fig. 2
As-received bloom microstructure: homogeneous bainite modified by tempering microscopy. Picral (a) and Nital (b,c,d)
etch. Optical microscopy at increasing magnifications (a,b,c) and electron microscopy (d).
Microstruttura del blumo allo stato di fornitura: bainite omogenea modificata dal rinvenimento. Attacco Picral (a) e Nital (b,c,d). Microscopia ottica ad ingrandimento crescente (a,b,c) e microscopi elettronica (d).
Moreover, some samples were subjected to the following laboratory re-heat-treatment: austenitization/solubilization at 1050 °C,
water quenching, double tempering at 400 °C.
Sets of either as-received or laboratory re-heat-treated samples
were then aged at three different temperatures: 470, 510, or 550
°C. Different samples of each set were extracted from the furnace after aging durations increasing up to 8 hours, and water
quenched.
The microstructure was examined by optical and electronic microscopy, after Nital or Picral [10] etch, and the austenitic average
grain size was measured by using the circular intercept method
[11], after Bechet-Beaujard [12] etch.
Standard tensile tests, plain-strain fracture toughness tests, Charpy-V impact tests, Vickers hardness tests, and FIMEC (Flat top cylindrical Indentations for Mechanical Characterization) test were
performed upon samples cut from the steel bloom, either in the
as-received state or after the above described re-heat treatments.
The reported hardness values are averages of 3 indentations.
Fracture toughness tests were performed on 35 mm thick SENB
(Single Edge Notch Bend) specimens [13]. The FIMEC indentation tests [14,15,16,17] were performed with a flat cylindrical indenter (1 mm diameter) and a 1.66 µm/s displacement rate. The
fracture surfaces of tensile and fracture toughness samples were
examined by Scanning Electron Microscopy (SEM).
X-ray diffraction and EDS analyses were performed on electrochemically extracted second phases (carbides and inclusions), in
order to detect the nature of the particles precipitated during the
aging heat treatment. The sample was dissolved in ethanol and
hydrochloric acid (10% vol.), the undissolved second phases were
collected on a filter (0.1 mm mesh size), and the filter was subjected to X-ray diffraction analysis (Co-Kα radiation). For comparison, the same analysis was carried out on an unused filter. EDS
analyses was performed on compacted second phase powder.
RESULTS
Microstructures
After the steelwork heat treatment, the as-received microstructure is homogeneous bainite, modified by tempering (Fig. 2).
Small randomly distributed carbide particles, not completely
resolved by optical microscopy, are present in the bainitic
matrix, probably Mo and V carbides. The previous austenite
grain boundaries are clearly evident (Fig. 2c,d), probably due
to the occurrence of a precipitated carbides layer, not always
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a
a
b
c
s
Fig. 3
Microstructures after aging for about 330 min
(from as-received condition) at 470 °C (a) and 550 °C (b);
temper modified bainite.
Microstrutture dopo invecchiamento per circa 330 min. (dalla
condizione di fornitura) a 470 °C (a) e 550 °C (b); bainite
modificata dal rinvenimento.
continuous, and partially removed during the metallographic
preparation. The average austenitic grain size is 130 µm (mean
of 263 intercepts).
As it regards the laboratory re-heat-treatment, after water
quenching, a martensitic microstructure with lath morphology
(typical of low carbon steels) was obtained. The microstructure
exhibits 445 HV100 hardness and 980 MPa yield strength (determined by FIMEC test). The tempered martensite, obtained after the first tempering at 400 °C, exhibited a morphology similar to the previous as-quenched martensite, but with a lower
hardness (420 HV100) and yield strength (910 MPa, by FIMEC
test). Both the microstructure and these mechanical properties
did not change noticeably after the second tempering.
The optical metallographic analysis of as received samples has not
detected important microstructural variation after aging (Fig. 3).
The previous austenite grain boundary precipitates, observed
in the as-received samples, are not present in the re-heat treated samples, neither after the quenching nor after the double
tempering at 400 °C (Fig. 4a); however, they are again visible,
in increasing amount, after aging the re-heat-treated samples
at temperatures increasing from 470 to 550 °C (Fig. 4b,c,d).
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s
Fig. 4
Re-heat-treated samples microstructure (water
quenching and double tempering at 400°C) (a). Microstructure after further aging for 330 min. at 470 °C (b),
510 °C (c) and 550 °C (d). Tempered martensite.
Microstruttura dei campioni ritrattati (tempra in acqua e
doppio rinvenimento a 400 °C) (a). Microstrutture dopo i
successivi trattamenti di invecchiamento per 330 min. a
470 °C (b), 510 °C (c) e 550 °C (d). Martensite rinvenuta.
Detailed SEM examination of aged samples (initially being either in the as-received or in the re-heat-treated condition) showed that the amount of detectable carbides mostly
decreases with the aging duration and temperature (Fig. 5);
therefore, it is hypothesized that during the aging treatments
the previously existing carbides, formed during the previous
heat treatments and probably not thermodynamically stable
in the aging temperature range, are progressively solubilized,
while finer (undetectable) carbides, possibly with a different
composition, are re-precipitated and may be the origin of the
detected hardness increase (see below).
The X-ray analysis of the second phases (carbides and inclusions) electrochemically extracted from the re-heat-treated
sample aged at 550 °C for 440 min is displayed in Fig. 6. The
most important diffraction peaks can be attributed to the
η-MoC carbide, with the possible presence of the V7C8 carbide. Moreover, the EDS analysis, carried out on the same
second phases compacted powder, has confirmed the occurrence of molybdenum, that has the highest peak, and of
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b
d
s
Fig. 6
Filter and filter-plus-carbides diffraction spectra
of re-heat treated sample aged at 550 °C for 440 min;
diffraction peaks of η-MoC (PDF # 08-0384) and V8C7
(PDF # 35-0786).
Spettri di diffrazione di un filtro vuoto e del filtro con
carburi del campione ri-trattato ed invecchiato a 550
°C per 440 min; picchi di diffrazione di η-MoC (PDF #
8-384) e V8C7 (PDF # 35-786).
s
Fig. 7
s
Fig. 5
Amount of detectable carbides (from SEM
observations) during aging of either as-received or
re-heat-treated samples, as a percentage of the
amount observed before aging, as a function of the
aging duration and temperature, for different carbide
morphologies (continuous lines - elongated carbides,
and dashed lines - small carbides).
Quantità di carburi rilevabili (da osservazioni SEM)
durante l’invecchiamento di campioni o in stato di
fornitura, o ri-trattati, come percentuale della quantità
osservata prima dell’invecchiamento, in funzione della
durata e temperatura dell’invecchiamento, per diverse
morfologie dei carburi (linee continue - carburi allungati e
linee tratteggiate - carburi piccoli).
Effect of the aging temperature and duration on the
hardness of the as received and re-heat-treated material.
Effetto della temperatura e della durata dell’invecchiamento
sulla durezza del materiale allo stato di fornitura e ri-trattato.
others elements, such as V, Fe, Cr, Si.
Mechanical tests
The results of the tensile and fracture toughness tests are listed
in Tab. 2 and 3, and compared with the previously assessed
properties of the ISO 1.2738 steel [2]. In particular, the fracture
toughness value of the examined steel in the as-received condition is somewhat higher than that of the ISO 1.2738 steel,
whereas the tensile properties are comparable.
The hardness curves relative to the age hardening heat treatment on the as-received and re-heat-treated samples are displayed in Fig. 7. The 550 °C aging temperature yielded the
5
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Condition / steel
As-received (A.r.)
A.r. + Aged 470 °C
A.r. + Aged 510 °C
A.r. + Aged 550 °C
RHT + aged 470 °C
RHT + aged 510 °C
RHT + aged 550 °C
1.2738 surface
1.2738 core
YS
MPa
855
817
1054
1225
1062
1154
1302
953
665
UTS
MPa
1156
1233
1331
1417
1288
1356
1457
1092
983
n
0.13
0.16
0.12
0.14
0.10
0.09
0.10
0.10
0.17
s
Tab. 2
Tensile properties of the examined steel, in different
metallurgical conditions, compared with the 1.2738 steel
(average values for each steel or position). YS: Yield Strength; UTS: Ultimate Tensile Strength; n: hardening exponent;
Elu uniform elongation; Elt elongation at fracture. Aging
duration 2.5 h.
Proprietà tensili dell’acciaio esaminato, in diverse condizioni
metallurgiche, confrontato con l’acciaio 1.2738 (valori medi per
ciascun acciaio o posizione). YS: tensione di snervamento; UTS:
tensione di rottura; n: esponente di incrudimento; Elu allungamento uniforme; Elt allungamento a rottura. Tempo di invecchiamento 2,5 h.
Condition / steel
As-received (A.r.)
A.r. + Aged (3 h at 525 °C)
1.2738 surface
1.2738 core
Fracture toughness
MPa√m
70
43
35
45
s
Tab. 3
<< Memorie
The FIMEC stress vs. displacement
curves, obtained on the as-received
and re-heat treated samples aged
for the longest durations (about 440
min, Fig. 8), confirm the above results: the yield stress (like the hardness) increases from 890 MPa in
the as-received condition, to 1230,
1275 and 1375 MPa, after aging at
470, 510 and 550 °C, respectively;
the as-received and re-heat treated
FIMEC curves become similar at
increasing aging temperature, and
those pertaining to the highest
aging temperature are consistent
with the slight overaging observed
in the re-heat treated samples (Fig. 7).
Charpy impact notch tests were carried out as a function of
temperature on the as-received and re-heat-treated samples,
the latter being either not aged, or aged at 520 °C for 2 h. (Fig.
9). The laboratory re-heat-treatment noticeably decreases the
brittle to ductile transition temperature, which anyway remains above the room temperature (Fig. 9). The fracture surfaces appearance, in the center of the samples, is always brittle
and consists mainly of cleavage zones.
The results of the tensile tests performed on the as-received
and re-heat treated specimens after aging at 470, 510 and 550
°C for 2.5 h are displayed in Tab. 2 and in Fig. 10. As expected,
aging at increasing temperatures yields higher yield and tensile strength, but lower uniform and fracture elongation. Consistently with the aforementioned hardness measurements
(Fig. 7), aged as-received and aged re-heat-treated specimens
have different uniform elongation, yield and tensile strength,
for the same aging temperature, but these differences decrease at increasing aging temperature. On the contrary, the
elongation-to-fracture difference increases; in particular, the
aged as-received samples fail in a brittle manner, without appreciable necking. The lower toughness of the as-received and
aged condition is confirmed by a significant reduction of the
fracture toughness after aging, from 70 to 43 MPa√m (Tab. 3).
The plane-strain fracture surface of the as-received material
Elu
%
8.7
9.2
6.3
2.4
5.2
4.6
2.8
7.0
8.8
Elt
%
15
12.1
7.6
1.1
13.6
12.9
12.0
14
15.7
Fracture toughness (KIc) of the examined steel, in
different metallurgical conditions, compared with the 1.2738
steel (average values).
Tenacità a frattura (KIc) dell’acciaio esaminato, in diverse condizioni metallurgiche, confrontato con l’acciaio 1.2738 (valori
medi).
highest hardness values: 490 HV100 for the re-heat treated
sample and 485 HV100 for the as-received one, starting from
420 and 380 HV100, respectively. As a consequence, the as-received specimens are more sensitive to the aging heat treatment than the re-heat-treated ones, achieving a similar hardness notwithstanding their hardness being less before aging.
Probably, this behavior is due to the fact that the hardness
of the bainitic and martensitic microstructure (of the as-received and re-heat treated samples, respectively) progressively becomes similar at increasing aging temperature, due
to supplementary tempering phenomena superimposing on
the aging precipitation. Overaging phenomena are not detected, except that in the re-heat treated sample aged at 550
°C, which exhibits a slight decrease in hardness for duration
longer than 3 hours.
6
s
Fig. 8
FIMEC flat indentation stress-displacement curves.
As-received and re-heat-treated (RHT) specimens before and
after aging for about 440 min at 470, 510, or 550 °C.
Curve tensione spostamento dell’indentazione piana FIMEC.
Campioni in stato di fornitura e ritrattati (RHT) prima e dopo
l’invecchiamento per circa 440 min a 470, 510 o 550 °C.
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s
Fig. 10
s
Fig. 9
Charpy-V impact tests: brittle-to-ductile transition
curves of the as-received, re-heat-treated, and re-heattreated and aged (at about 520°C for 2 h) metallurgical
conditions.
Prove di resilienza Charpy-V: curve di transizione fragile-duttile
dei campioni in condizioni metallurgiche di fornitura, ri-trattata,
e ri-trattata ed invecchiata (a circa 520 °C per 2h).
shows mainly cleavage facets (Fig. 11a,b), with small ductile
intergranular rupture areas (Fig. 11b); the latter morphology
becomes prevalent in the as-received tensile fracture surface,
together with some cleavage (Fig. 12).
Overall, the morphology of the tensile fracture surfaces of the
different examined aged conditions depend mainly on the
metallurgical state before the aging heat treatment. In fact,
both the aged and not-aged as-received samples exhibit cleavage areas and ductile intergranular rupture areas (consistently
a
Tensile properties of aged samples (from initial asreceived and re-heat treated, RHT, condition). Yield Strength
(YS) and Ultimate Tensile Strength (UTS), elongation to
fracture (Elt) and uniform elongation (Elu).
Proprietà tensili di campioni invecchiati (dalle condizioni iniziali
di fornitura, as. rec., o di ri-trattamento, RHT). Tensione di
snervamento (YS) e di rottura (UTS), allungamento a rottura
(Elt) ed allungamento uniforme (Elu).
with the lack of necking, Fig. 12a,c,e), whereas the fracture
surfaces of the re-heat-treated and aged samples always show
a cup-and-cone morphology, with mode-I coalesced microvoids and mode-II shear areas (Fig. 12b,d).
DISCUSSION AND CONCLUSIONS
The microstructure of the examined positions inside the steel
bloom consists almost completely of bainite modified by tempering. Therefore, the bloom fracture toughness is low in comparison to usual quenched and tempered steels, being about
b
s
Fig. 11
Plane-strain fracture surfaces in the as-received steel, at the onset of metastable crack propagation (a) and in the crack
propagation region (b).
Superfici di frattura in deformazione piana nell’acciaio in stato di fornitura, all’inizio della propagazione instabile (a) e nella regione di
propagazione instabile (b).
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b
c
d
d
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Fig. 12
Tensile fracture surface of as-received (a,c,e) and
re-heat-treated (b,d) steel, after aging at 550 °C. Overviews
(a,b) and details (c,d,e).
Superfici di frattura a trazione dell’acciaio in stato di fornitura
(a,c,e) e ri-trattato (b,d), dopo invecchiamento a 550 °C. Viste
complessive (a,b) e dettagli (c,d,e).
70 MPa√m, but somewhat higher in respect to the largely
used 1.2738 steel (40 MPa√m on average).
The as-received steel shows a precipitates layer at the previous austenitic grain boundaries, which may be tentatively
related with the ductile intergranular fracture observed after the tensile tests in the as-received, both not-aged and
aged, samples. Nevertheless, the aged re-heat-treated tensile specimens present fully ductile fracture surfaces, even
if they also show a similar precipitates layer at the austenitic grain boundaries after the aging treatment (although
not after the re-heat-treatment). Therefore, the possible relationship among the grain-boundary precipitates and the
intergranular rupture is not yet completely clear.
The as-received specimens are more sensitive to the aging
heat treatment than the re-heat-treated ones, since they
yield increasingly similar hardness values after aging at increasing temperatures, notwithstanding their lower hardness before aging. This fact may be partially explained by
hypothesizing that the initial differences between the bainitic and martensitic matrixes (already tempered at low
temperature) are progressively reduced due to coalescence
phenomena occurring during tempering at increasing temperatures and interferring with the precipitation during aging.
The FIMEC and tensile tests of aged and not aged samples
overall confirm the results of the hardness versus aging duration tests; moreover, the tensile tests evidence a difference
in the fracture mode between the as-received and aged samples and those re-heat-treated and aged, the former generally showing brittle fracture surfaces without significant
necking, and the latter ductile fracture surfaces and evident
necking. Therefore, and also by considering the reduction
of fracture toughness (43 MPa√m) of the as-received steel
after aging at 525 °C, it is concluded that the aging treatment generally causes a relevant toughness reduction.
Overall, the reported results, and particularly the hardness
curves as a function of the aging duration, outline the kinetics of the aging process and constitute a data set that could
be employed for the choice of the more suitable parameters
(duration and temperature) for the aging treatment of specific molds. In particular, the substantially asymptotic trend
of the hardness, as a function of the aging duration, at the
examined temperatures and for durations up to 8 h, may
allow to obtain homogeneous result also in the aging of
molds with large cross-section, for which the actual duration at temperature is necessarily differentiated from surface to core.
Nevertheless, the nature of the metallurgical transformations that originate the hardening process has not been
completely determined yet and will be the subject of further
studies. A first hint in this direction is given by the observation of the gradual disappearance of the previously existing
carbides during the aging treatments, which may be related
with the precipitation of different, finer carbides.
ACKNOWLEDGEMENTS
Italian Ministry for University and Research, for financial support by research grant PRIN 2005090102. Lucchini Sidermeccanica steelwork, Lovere, Italy, for steel procurement and the
CCT diagram.
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2001 steel, in: Proceedings of 7th international tooling conference - Tooling materials and their applications from research to
market - Politecnico di Torino – Torino, Italy, 2-5 may 2006, curatori M. Rosso, M. Actis Grande, D. Ugues (Milano: Ancora,
2006), Vol. 2, 153-161.
10] E407-99, Standard Practice for Microetching Metals and Alloys”, ASTM, 1999.
11] E112-96. Standard test methods for determining average
grain size. ASTM, 1996.
12] S. Bechet, L. Beaujard. Nouveau réactif pour la mise en évidence micrographique du grain austénitique des aciers trempés ou trempés - revenus,” La Revue de Métallurgie, 52 (1955),
830.
13] E399-05. Standard test method for plane-strain fracture
toughness of metallic materials. ASTM, 2005.
14] A. Donato, P. Gondi, R. Montanari, L. Moreschi, A. Sili, S.
Storai. A remotely operated FIMEC apparatus for the mechanical characterization of neutron irradiated materials. Journal of
Materials Science, 258-263 (1998), 446-451.
15] R. Mouginot, D. Maugis. Fracture indentation beneath flat
and spherical punches, Journal of Materials Science, 20 (1985),
4354-4376.
16] H.Y. Yu, M.A. Imam, B.B. Rath. Study of the deformation
behaviour of homogeneous materials by impression tests,
Journal of Materials Science, 20 (1985), 636-642.
17] P. Gondi, R. Montanari, A. Sili. Small-scale nondestructive
stress-strain and creep tests feasible during irradiation, Journal
of Nuclear Materials, 212-215 (1994), 1688-1692.
9
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<< Memorie
ABSTRACT
EFFETTO DEL TRATTAMENTO TERMICO SULLE
PROPRIETÀ MECCANICHE DI UN ACCIAIO INDURENTE
PER PRECIPITAZIONE PER GRANDI STAMPI PER
MATERIE PLASTICHE
Parole chiave: acciaio, precipitazione
Gli stampi per particolari in materia plastica di grande dimensione, quali
per esempio paraurti e cruscotti, usati nell’industria automobilistica, sono
solitamente lavorati per asportazione di truciolo da grandi blocchi di acciaio prebonificato. I blumi di acciaio ISO 1.2738 tradizionalmente usati presentano microstrutture disomogenee e tenacità ridotta (KIc ≈ 40 MPa√m
e KV ≈ 10 J a temperatura ambiente); inoltre questo acciaio è difficilmente
saldabile (Ceq ≈ 1,16). Numerosi acciai indurenti per precipitazione sono
stati proposti come alternativa per superare questi limiti.
Siccome i processi di indurimento per precipitazione inducono deformazioni molto limitate, possono essere svolti dopo la lavorazione meccanica.
Inoltre, questi processi possono produrre microstrutture e proprietà meccaniche omogenee in stampi grandi e geometricamente complessi.
È stato esaminato un acciaio induribile per precipitazione sviluppato per la
fabbricazione di stampi per materie plastiche, con la composizione chimica
esposta in Tabella I. Il ciclo produttivo consiste di colata, forgiatura, trattamento termico iniziale (svolto in acciaieria), lavorazione meccanica ed
indurimento per invecchiamento. Il trattamento termico iniziale consiste
di un trattamento preliminare (includente una de-idrogenazione) seguito
da austenitizzazione/solubilizzazione a 1020 °C, tempra in aria e doppio
rinvenimento a 400 °C (Fig. 1); questo trattamento è rivolto ad ottenere
una microstruttura bainitica con durezza compresa tra 310 e 350 HB.
Un blumo di dimensione originale 500x1500x2400 mm è stato esaminato
allo stato di fornitura, cioè dopo il trattamento termico iniziale sopra descritto. Inoltre, alcuni campioni sono stati ritrattati in laboratorio come
segue: austenitizzazione/solubilizzazione a 1050 °C, tempra in acqua, doppio rinvenimento a 400 °C, ottenendo martensite rinvenuta.
Il rafforzamento per precipitazione è stato studiato usando campioni ini-
10
zialmente o allo stato di fornitura, o ritrattati in laboratorio, e poi invecchiati a 470, 510 o 550 °C per durate fino ad 8 h.
Le microstrutture sono illustrate nelle Fig. 2, 3 e 4. La dimensione del
grano austenitico allo stato di fornitura è di 130 µm. Carburi precipitati
presso i bordi di grano austenitici sono presenti allo stato di fornitura,
assenti allo stato ri-trattato, e nuovamente presenti allo stato ritrattato
ed invecchiato. La quantità di carburi osservati mediante SEM decresce
all’aumentare della temperatura e durata dell’invecchiamento (Fig. 5); si
ipotizza che i carburi esistenti agli inizi degli invecchiamenti siano gradualmente solubilizzati, permettendo la precipitazione di altri carburi più
fini, non osservabili, che causano il rafforzamento. L’analisi XRD dalle
seconde fasi, estratte per via elettrochimica (da un campione ri-trattato ed
invecchiato a 550 °C, Fig. 6), evidenzia la presenza di η-MoC e la possibile
presenza di V7C8.
Allo stato di fornitura la durezza e le proprietà tensili sono paragonabili a
quelle dell’acciaio ISO 1.2738, ma la tenacità a frattura è maggiore (Tab. II
e III); la frattura in deformazione piana avviene per clivaggio (Fig. 11). Il
ritrattamento termico riduce sensibilmente la temperatura di transizione
fragile-duttile (Fig. 9). Gli invecchiamenti incrementano notevolmente la
durezza (Fig. 7); l’incremento è maggiore per i campioni allo stato di fornitura, la cui durezza iniziale è minore; per le temperature e durate esaminate il sovrainvecchiamento è assente o trascurabile. Questi risultati sono
confermati anche da prove di indentazione strumentata con penetratore
cilindrico (FIMEC, Fig. 8). L’invecchiamento a temperature crescenti aumenta le tensioni di snervamento e rottura, coerentemente con l’aumento
della durezza (Fig. 10); però i campioni di trazione allo stato di fornitura ed invecchiati si rompono in modo fragile, senza apprezzabile strizione
(contrariamente a quelli ritrattati ed invecchiati, Fig. 10 e 12). L’effetto
avverso dell’invecchiamento sulla tenacità, nel caso di materiale inizialmente allo stato di fornitura, è confermato dalla riduzione della tenacità a
frattura (KIc ≈ 43 MPa√m dopo invecchiamento a 525 °C).
Nel loro insieme, i risultati esposti illustrano la cinetica e gli effetti del
processo di invecchiamento e costituiscono un insieme di dati utili per scegliere i parametri di invecchiamento più adeguati per specifici stampi.
Memorie >>
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HOW HEAT TREATMENT CAN GIVE
BETTER PROPERTIES TO ELECTROLESS
NICKEL-BORON COATINGS
V. Vitry, F. Delaunois, C. Dumortier
Electroless nickel-boron deposits were synthesized on mild steel and submitted to heat treatments under neutral and nitrogen based atmosphere. The properties obtained after these treatments were compared to as deposited nickel-boron coatings. The morphology and structure of the samples were investigated by XRD, SEM
and optical microscopy; their composition was studied by ICP, GD-OES and SIMS analysis, and micro and
nanoindentation tests were carried out to assess the coatings’ hardness. Scratch tests were used to determine
the damage mechanisms of the coating.
KEYWORDS: electroless deposition – nickel-boron – nanoindentation – heat treatment
INTRODUCTION
Autocatalytic (Electroless) nickel plating was discovered by
Brenner and Riddel in 1946 [1]. This process is based on the
aqueous reduction of nickel salts by a chemical agent thus allowing deposition on non-conducting materials and leading to
continuous coatings with a constant thickness [2-5]. Nickel boron coatings are obtained when a boron-based agent, such as
sodium borohydride is used to reduce the nickel. Those coatings are of great interest and are extensively studied [6-12].
They present, in their as-deposited state, an hardness close to
750hv100 and are useful in many industries including automotive, electronic and chemical industries because of their good
mechanical, chemical and tribological properties [2, 5, 11, 13,
14]. Depending on the amount of boron present, the coatings
are considered amorphous, microcrystalline or a mix of the two,
the amount of amorphous phases increasing with the amount of
boron [2, 7, 15-18].
Heat treatments are often used to enhance the properties of
nickel-boron coatings: they allow crystallisation of the amorphous part and, if well designed, lead to nano and microcrystalline structure which are harder than the as-deposited coatings
and their hardness can reach 1200hv100. [2, 17, 19]
Much information can be obtained using nanoindentation: this
technique is an instrumented indentation and the loading and
unloading curves are recorded during each indent. Moreover,
the loads are much smaller than in the case of microindentation (typically a few mN) [20-22]. This technique is often used
with a Berkovitch indenter which has the same surface than the
Vickers indenter while being easier to manufacture owing its
Véronique Vitry, Fabienne Delaunois, Christian Dumortier
Mons Faculty of Engineering, Metallurgy Department,
Mons, Belgium
triangle-based pyramid shape. Working with very low loads allows to get very small indents and thus to study the hardness
evolution across a relatively thin coating.
Scratch test [23-29] can give information about the “practical
adhesion” of coatings as well as the degradation modes of the
coatings. It consists in the application of an increasing load to a
coating. Modern investigation techniques are used to study the
coating’s scratch test comportment: acoustic emission, friction
coefficient and penetration depth measurements are recorded
during the test and microscopic examination is carried out after
the test. The critical load of a system which characterizes the adhesive strength of the coating/substrate system is determined
from the first adhesive failure. The degradation modes can be
identified from observation. However scratch tests cannot be
used to predict quantitative wear rates of materials and coatings.
EXPERIMENTAL
Samples preparation
Steel and Aluminium alloy cylinders with a diameter of 25 ± 1
mm and a thickness of 10 ± 1 mm were plated with nickel-boron.
Before plating, they were mechanically polished, degreased with
acetone and etched in an acid solution. The aluminium samples
were subjected to further pre-treatment by double-zincate conversion and acid nickel phosphorous flash deposition.
The deposition bath is based on the reduction by sodium
borohydride (NaBH4); the nickel ions source is nickel chloride
(NiCl2.6H2O). The nickel ions are complexed by ethylene diamine (EN) and lead tungstate (PbWO4) is used as a stabilizer.
The operation conditions and the installation have been described elsewhere [15].
Classical heat treatments were carried out under neutral gas
flow (95%Ar – 5%H2) at 400°C for 1 hour for steel substrates and
at 180°C for 4 hours for aluminium substrates (this temperature
was proven by Delaunois et al. to offer a good compromise bela metallurgia italiana >> aprile 2009
1
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tween coating hardening and substrate softening for aluminium
substrates [15]). Preliminary test for other treatments were carried out under ammonia and nitrogen-based atmospheres.
Samples prepared for SEM observation and hardness testing
were first cut using a Leco Microtom cutting machine with a
diamond cutting disk, then mounted in a non-retractable resin
and mirror polished.
s
Fig. 1
GD-OES depth profile of an untreated nickelboron coating on Steel.
Profilo GD-OES nello spessore di un rivestimento nichelboro non trattato su acciaio.
s
Fig. 2
GD-OES depth profile of a nickel-boron coating
on Steel after 1 hour heat treatment at 400°C under
neutral atmosphere.
Profilo GD-OES nello spessore di un rivestimento nichelboro su acciaio dopo 1 ora di trattamento termico a
400°C in atmosfera inerte.
s
Fig. 3
SEM micrograph of an untreated nickel-boron
coating – left: cross section showing the columnar
morphology – right: cauliflower-like surface.
Micrografia al SEM di un rivestimento nichel-boro non
trattato – sinistra: sezione trasversale che mostra una
morfologia colonnare – destra: superficie “a cavolfiore”.
2
Samples analysis
Mean chemistry of the samples was investigated after dissolution in concentrated nitric acid by ICP analysis using a JobinYvon apparatus.
Profile analysis was carried out by GD-OES using Jobin-Yvon
spectrometer and surface analysis by TOF-SIMS using an IONTOF IV apparatus.
A Siemens D500 X-rays θ–θ apparatus applying Cu Kα (1,54 Ǻ)
radiation was used to study the structure of the samples. Their
morphology and thickness were observed using a Philips XL 20
Scanning Electron Microscope.
Microhardness measurements were carried out using a LECO
M-400-A, mounted with a Vickers indenter for surface testing
and with a Knoop (lozenge-shaped) indenter for cross section
testing. A load of 100g was used for Vickers indentation while a
load of 50g was used for Knoop hardness testing. The holding
time was 20s for both techniques.
Nanohardness was obtained with a MTS nano-indenter XP
mounted with a Berkovitch (tetrahedron shaped), using depth
controlled indentation in order to obtain indents of similar size.
The hardness value at a load of 4000μN was chosen as nanoindentation hardness value.
Scratch tests were performed on selected samples using the continuous load increase method up to 30N with a Microphotonics
Micro Scratch Tester (MST), with a load rate of 19.17N/min and
an advance rate of 9.58mm/min, resulting in a scratch of 15mm.
The tip was a Rockwell C diamond stylus indenter with a radius
of 200 μm.
RESULTS AND DISCUSSION
Chemistry of deposits
ICP analysis of as-deposited coatings showed that their average
composition is 93 wt. % nickel, 6 wt. % boron and 1 wt. % lead.
Heat treatments do not modify the global composition of the
coatings.
GD-OES analysis, on steel substrate, allowed us to follow qualitatively the composition into the depth of the coating. In the asplated state (Fig. 1), the boron and nickel content of the coating
don’t vary with the depth while the lead content seems to be
higher at near the substrate interface then decreases slightly before increasing once more. It means that more lead is deposited
at the very beginning of the process and at the end that during
the fast “regime” deposition of the coating.
Observation of the interface allows us to predict a good adhesion of the coating because there seems to be a certain amount
of interdiffusion with the substrate.
After heat treatment at 400°C for one hour under neutral atmosphere (Fig. 2), the nickel and boron content are quite unmodified
while the lead content becomes higher near the free surface of
the sample and decreases steeply with depth. The lead seems to
diffuse outside the coating. However, the interface between the
steel substrate and nickel-boron was not modified by this treatment. Previous work revealed important interdiffusion at the
interface after heat treatment in the case of aluminium-silicon
alloys (AS7G06) [15].
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Knoop microhardness (hk50)
Vickers microhardness (hv100)
Berkovitch nanoindentation (4000μN)
Untreated
834 ±20
854 ± 40
823± 155
4h ; 180°C
927 ± 30
1014 ± 40
1140 ± 75
1h ; 400°C
1302 ± 40
1584 ± 182
s
Tab. 1
Vickers and Knoop and Berkovitch hardness values of nickel-boron coatings on aluminium alloys.
Valori di durezza Vickers, Knoop e Berkovitch di rivestimenti di nichel-boro su leghe di alluminio.
Positive and negative ions SIMS analysis was carried out on the
untreated samples and revealed only the presence of the known
constituent of the coating (Nickel, boron, lead) and of the classic
surface contamination. The results obtained after neutral atmosphere heat treatment were similar, proving that the coating’s
chemistry is not much influenced by those treatments.
Structure and morphology of the coatings
In the as-deposited state, the coatings present a columnar morphology and a cauliflower-like surface (which is characteristic
of nickel-born coatings [11,15,17] ), as can be seen on Fig. 3.
Neutral atmosphere heat treatments up to 400°C do not modify
those properties.
The structure of untreated samples and samples treated at 180°C
revealed they were amorphous (Fig. 4) while crystallization occurred during heat treatment at 400°C. This is expected from the
literature and our own previous results [2,15,19,30].
The effect of this crystallization on the mechanical properties of
the coating will be discussed later.
Mechanical properties of the coating
Vickers hardness testing on the unprocessed surface of the sample is the standard method to measure hardness of nickel-boron
coatings. However, we find it disputable because the surface is
unprepared and its smoothness is unwarranted, and because
the substrate hardness may significantly influence the results
when the applied load is too high. We thus used other hardness
testing methods to free ourselves from those potential problems:
Knoop microindentation and Berkovitch nanoindentation were
used on polished cross-sections. Nanoindentation were converted from GPa into Berkovitch hardness points (equivalent to
Vickers values) to facilitate comparison.
Hardness values for untreated samples were close to 825 for
all techniques. After heat treatment at 180°C for 4 hours, those
values reached 1000, and they were further increased after heat
treatment at 400°C. The first increase is caused by short order rearrangement in the coating (the amorphous dome XRD intensity is slightly higher), while the crystallization observed between
180°C and 400°C causes a far greater hardness enhancement and
reaches the maximum hardness value for nickel-boron coatings
[2,17,19]. It is due to the generation of a high density of grain
boundaries inside the coating. The hardness can thus be optimized by the grain-size control: if the grains are allowed to coalesce (i.e. if the heat treatment is too long or the temperature too
high), the hardness of the coating will decrease [16.17,31].
It was not possible to obtain a reliable Knoop hardness value after treatment at 400°C because cracking of the coating occurred.
This shows the importance of comparing values from different techniques: Nickel boron coatings have a very anisotropic
structure, due to their synthesis mode. Vickers hardness is carried out in the growth direction of the coating and any damage occurring during this test will remain unseen because it will
take place inside the coating. However, Knoop indentation is
made perpendicular to the growth direction of the coating and
the subsequent damage is easily observable (Fig. 5). Knoop indentation is thus more reliable. This may also explain why na-
s
Fig. 4
X-Ray diffraction patterns of Electroless nickelboron coatings on aluminium substrates with and
without heat treatments.
Spettri di diffrazione ai raggi X per i rivestimenti chimici di
nichel-boro su substrati di alluminio con e senza trattamenti termici.
s
Fig. 5
Cracking of the coating caused by Knoop
indentation.
Criccatura del rivestimento causata dall’impronta Knoop.
3
Trattamenti termici
s
Fig. 6
Nanoindentation profiles on treated and
untreated nickel-boron coatings.
Profili di nanodurezza su rivestimenti nichel-boro trattati
e non trattati.
noindentation values are so much higher than Vickers values for
400°C treatment: the load applied for the nanoindentation test is
too low to generate damage.
The small size of nanoindentation indents allows measuring
the hardness of the coating on several points across the coating without any interaction of the indents. We were thus able to
draw hardness profiles for nickel-boron coatings (fig. 6). Those
profiles showed that the hardness values are not homogeneous
across the coating except after heat treatment at 180°C. This is
probably due to the fact that as-deposited coatings have important internal stress. Heat treatment at 180°C releases a lot of this
stress and the hardness becomes homogeneous. Heat treatment
at 400°C however may generate new internal stresses because it
is accompanied by important structural modifications. Moreover, as nanoindentation is a very sensitive experimental method
it is possible hat part of the values scattering is due to external
factors such as roughness of the sample (microscratches in the
s
Fig. 7
Micrograph of scratches on 6μm thick nickelboron coatings (aluminium alloy substrate) without
heat treatment.
Micrografia di scalfitture su rivestimenti di nichel-boro
dello spessore di 6μm (substrato in lega di alluminio)
senza trattamento termico.
4
<< Memorie
polished cross section), geometry of the indenter tip, thermal
drift and creep.
The usual thickness used for nickel-boron coatings is at least 1525 micrometers. For coatings of this thickness, scratch tests up
to 25N did not reveal any significant damage. We decided thus
to work with thinner coatings (with a thickness of 6µm) in order
to determine the damage mechanisms observed during scratch
test. These tests were only carried out without heat treatment
and after heat treatment at 180°C.
The damage mechanisms were the same for both treated (Fig. 8)
and untreated (Fig. 7) coatings which is not surprising knowing
they present roughly the same structure and morphology. The
first failure that was observed is longitudinal cracking on the
edges of the coating. Chevron cracks were then observed. Those
are cohesive damage. The first adhesive failure is discontinuous
ductile perforation of the coating, which is followed, in the case
of heat treated samples, by continuous ductile perforation of the
coating. To observe this kind of failure is really encouraging because it is a proof of the good adhesion of the coating, which
was first suggested by the interface observation by GD-OES.
Moreover, this failure was only observed because the coating
was so thin.
Effects of alternative post-treatment
Alternative heat treatments are now in their testing phase. They
are mainly based on the use of reactive atmospheres. We present
hereunder some interesting results we obtained using (i) a treatment under a reduced pressure in a nitrogen-based gas, which
will be called “vacuum treatment” in further explanations and
(ii) a treatment under an ammonia-based atmosphere, which
will be called “ ammonia treatment”.
SEM cross section micrographs of nickel-boron coatings after
“vacuum” and “ammonia” treatments showed an important
morphological modification of the coating: it becomes totally
dense after “vacuum” treatment and the columnar morphology
disappears completely (Fig. 9a), and the coating is composed
s
Fig. 8
Micrograph of scratches on 6μm thick nickelboron coatings (aluminium alloy substrate) after 4
hours at 180°C under 95% Ar + 5% H2 atmosphere.
Micrografia di scalfitture su rivestimenti di nichel-boro
dello spessore di 6μm (substrato in lega di alluminio)
dopo 4 ore a 180°C in atmosfera 95% Ar + 5% H2.
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Trattamenti termici
of 2 distinct layers after “ammonia” treatment (Fig. 9b). The inner layer is dense and resembles the “vacuum” treated coating
while the outer layer looks porous.
Knoop hardness measurements were carried out on “vacuum”
treated samples and on the dense part of “ammonia” treated
samples. Values of 1570 ± 100 hk25 and 1630 ± 100 hk25 respectively were obtained instead of 1250 ± 100hk25 after treatment at
400°C. It shows that the hardness of nickel-boron coatings can
be further enhanced by the use of modified heat treatments.
Scratch tests were carried out on “ammonia” treated coatings.
Those coatings, while they are plastically deformed, are nearly
undamaged after the tests (Fig. 10) which is promising for wear
applications.
CONCLUSIONS
As-deposited nickel-boron coatings possess several interesting
features: high hardness (~825hv100), good scratch comportment,
amorphous structure, columnar morphology and cauliflowerlike surface.
Heat treatment influences some of those features, mostly in a
positive way:
- The columnar morphology is unmodified by classical heat
treatment up to 400°C but is transformed in a dense layer that
can be accompanied by a porous outer layer after the alternative
treatments we investigated.
- The amorphous structure undergoes crystallization during
heat treatment if the temperature is high enough: there is no
crystallization for 180°C treatments while crystallization is complete after treatment at 400°C.
- The hardness of the coating is very much influenced by its
crystalline state: while low temperature treatment induces a
slight increase and an homogenisation of the hardness, the treatment at 400°C leads to an hardness value of 1300hv100. This high
values is due to the important grain-boundaries density that is
obtained after heat treatment.
- Alternative heat treatment allowed a further hardness increase
by a still unidentified mechanism.
- The scratch comportment of nickel-boron coatings is quite
unmodified by neutral heat treatments. The comportment after
ammonia-based alternative treatment is mainly plastic deformation of the outer layer of the coating, which is very interesting
for wear applications.
- Adhesion of the coating is predicted to be good because of
the chemical interaction seen at the coating/substrate interface
and of the scratch comportment. Heat treatments don’t seem to
modify the interface.
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s
Fig. 9
SEM micrograph of (a) a “vacuum” treated and
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Micrografia al SEM di una deposizione di nichel-boro su
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s
Fig. 10
Scratch test on an “ammonia” treated nickelboron coating.
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5
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<< Memorie
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29] M. Čekada, J. Dolinšek, P. Panjan, Vacuum, 80, (1-3), (2005)
137-140.
30] V. Vitry, F. Delaunois and C. Dumortier, Surface and Coatings Technology, 202, 14 (2008) 3316 - 3324.
31] B. Oraon, G. Majumdar, B. Ghosh, Materials and Design
(2007) vol 29 (7) 1412 - 1418.
ABSTRACT
COME IL TRATTAMENTO TERMICO PUÒ
MIGLIORARE LE PROPRIETÀ DEI RIVESTIMENTI
CHIMICI DI NICHEL-BORO.
Parole chiave: trattamenti termici, rivestimenti, acciaio
Deposizioni chimiche di nichel-boro sono state realizzate su un acciaio
dolce e sono state poi sottoposte a trattamenti termici in atmosfera inerte
6
o a base di azoto. Le proprietà ottenute dopo questi trattamenti sono state
confrontate con quelle dei rivestimenti nichel-boro di partenza. La morfologia e la struttura dei campioni sono state esaminate mediante XRD,
microscopia ottica e SEM; la loro composizione è stata studiata mediante
analisi ICP, GD-OES e SIMS; infine sono state eseguite prove di micro- e
nano-durezza per valutare la durezza dei rivestimenti. Per determinare i
meccanismi di danneggiamento del rivestimento sono stati utilizzati prove di resistenza alla scalfittura.
Memorie >>
Trattamenti termici
IMPIEGO DI PROTOSSIDO D’AZOTO
NEL TRATTAMENTO TERMOCHIMICO
DI NITRURAZIONE GASSOSA:
STUDIO DEI PROCESSI PRODUTTIVI E
CARATTERIZZAZIONE METALLURGICA
M. Merlin, P. Camanzi, G. L. Garagnani
In questo lavoro sono presentati i risultati di uno studio effettuato sul trattamento termochimico di
nitrurazione gassosa, in particolare sulla possibilità di ottenere un aumento della velocità di processo a livello
industriale. È stato testato l’effetto dovuto all’introduzione, insieme all’ammoniaca anidra (NH3), di protossido
di azoto (N2O) nell’atmosfera di processo. Inizialmente sono stati caratterizzati i normali cicli produttivi
di nitrurazione gassosa in atmosfera di ammoniaca anidra, inserendo campioni identici di acciaio di tipo
42CrMo4 all’interno delle cariche di materiale da nitrurare nei forni a pozzo. I cicli produttivi di interesse
sono stati quelli aventi come obiettivo finale l’ottenimento di uno strato di indurimento superficiale di 3 e 4
decimi di millimetro. Al termine della sperimentazione si è constatato che l’utilizzo dell’N2O comporta un
incremento della cinetica di processo, con una conseguente riduzione del 20÷30% dei tempi di processo rispetto
al ciclo produttivo standard. Nonostante la riduzione dei tempi di processo, le caratteristiche meccanicomicrostrutturali dei campioni trattati sono risultate confrontabili con il processo tradizionale, prospettando
l’idoneità all’applicazione industriale della nitrurazione gassosa con impiego di protossido d’azoto. Alla
sperimentazione e raccolta dei dati ottenuti segue una valutazione economica e di fattibilità per i cicli
produttivi testati, sulla base di un confronto qualitativo con i processi già applicati industrialmente.
PAROLE CHIAVE: protossido d’azoto, indurimento superficiale, metallografia, nitrurazione gassosa
INTRODUZIONE
Le odierne esigenze produttive nel campo della meccanica
portano a richieste sempre maggiori di elevate caratteristiche
tecniche degli organi meccanici progettati. L’obiettivo infatti
è quello di consentire a tali particolari un impiego in condizioni di esercizio particolarmente onerose in termini di tensioni meccaniche, di usura o di corrosione. Spesso la tecnologia
produttiva che porta all’ottenimento dei particolari meccanici
non consente di lavorare agevolmente metalli con determinate
caratteristiche; basti pensare, ad esempio, alle lavorazioni per
asportazione di truciolo eseguite su materiali ad elevata durezza. Inoltre, sono frequenti i casi in cui le caratteristiche richieste per l’utilizzo non sono nemmeno riscontrabili nei materiaMattia Merlin, Gian Luca Garagnani
Dipartimento di Ingegneria, Università degli Studi di Ferrara
Paolo Camanzi
Siderit s.r.l., Zola Predosa (Bologna)
li di maggiore utilizzo, alle condizioni di normale fornitura.
Pertanto si rende spesso necessaria l’esecuzione di trattamenti
termici e termochimici per conferire al componente le dovute
proprietà meccaniche. In particolare, i trattamenti termochimici prevedono l’apporto di uno o più elementi chimici dall’esterno, in modo tale da alterare superficialmente la composizione
chimica, oltre che quella strutturale, del materiale iniziale.
La nitrurazione gassosa è un trattamento termochimico di
indurimento superficiale che viene condotto a temperature
comprese tra i 480°C e i 570°C per un periodo di tempo variabile da qualche ora a parecchie decine di ore, in condizioni tali da consentire una diffusione di azoto nella superficie
del particolare da trattare. Al termine del trattamento il pezzo
viene raffreddato a temperatura ambiente mediante un’atmosfera gassosa di tipo inerte. Gli strati superficiali così generati
sono caratterizzati da un’ottima resistenza all’usura adesiva,
al grippaggio e all’abrasione meccanica. La nitrurazione può
essere eseguita su numerosi tipi di acciaio, anche se le caratteristiche di durezza superficiale sono massime per gli acciai legati contenenti alluminio, cromo e vanadio. I più diffusi acciai
1
Trattamenti termici
C
Simax
Mn
42 CrMo4 0.38-0.45 0.4 0.60-0.90
Pmax
0.025
Smax
0.030
<< Memorie
Al
-
Cr
0.90-1.20
Mo
0.15-0.25
Ni
≤ 0.30
V
-
s
Tab. 1
Composizione nominale (percentuale in peso) dell’acciaio 42CrMo4.
Chemical composition (range of wt%) of the 42CrMo4 steel.
IMPIANTO 1
IMPIANTO 2
Tipo
Forno a pozzo
Forno a pozzo
Diametro [cm]
150
120
s
Tab. 2
Impianti produttivi.
Productive implants.
da nitrurazione in commercio sono, secondo la Norma UNI
EN 10085, il 31CrMo12 e il 41CrAlMo7, a cui si aggiungono
anche il 31CrMoV9 e il 34CrAlMo5. Non sempre durezze eccessivamente elevate sono desiderabili per impieghi costruttivi e pertanto il trattamento di nitrurazione gassosa può essere
eseguito anche su acciai più basso legati [1-9]. Un esempio tipico è costituito dall’acciaio da bonifica basso legato 42CrMo4,
che viene normalmente utilizzato per la realizzazione di molti
organi meccanici in quanto presenta buone proprietà di tenacità e resistenza. Su tali componenti in 42CrMo4 spesso viene
richiesto un trattamento superficiale di nitrurazione gassosa,
pur non essendo un acciaio tipicamente “da nitrurazione”.
I normali cicli di nitrurazione gassosa per elevata profondità
di indurimento comportano tempi di impegno impianto molto
elevati. Pertanto, è stata valutata la possibilità di incrementare
la velocità di processo tramite l’impiego di uno specifico gas
di processo addizionale: il protossido di azoto (N2O). L’effetto
dato dal protossido d’azoto si traduce in una maggiore capacità nitrurante dell’atmosfera e quindi in un abbassamento dei
tempi di processo. I risultati sperimentali di molteplici autori
[10-24] mettono in evidenza questi effetti. In particolare, da
[17, 23-24] si può concludere che una addizione di ossigeno
all’atmosfera nitrurante porta ad una azione benefica, specialmente per acciai legati e contenenti cromo. Si è incominciato a
sfruttare questi principi già dalla metà degli anni ottanta per
velocizzare i trattamenti di nitrurazione a livello industriale; la
società Air Liquide ha sviluppato e depositato brevetti per processi denominati ALNAT-N®, ALNAT-NO®, i quali prevedono
l’impiego di protossido d’azoto come gas di processo [25-26].
Un impiego non controllato di protossido d’azoto, specialmente nella fase iniziale del processo di nitrurazione, può portare
a difetti sul materiale trattato; infatti, l’aumento delle caratteristiche meccaniche conferito da questo trattamento può talvolta
accompagnarsi ad alcuni inconvenienti. Una iper-nitrurazione
può generare una coltre bianca di eccessivo spessore, con conseguente fragilità della stessa. Un altro difetto riguarda l’eccessiva precipitazione di nitruri entro lo strato di diffusione;
è buona norma evitare la formazione di reticoli continui di
nitruri aciculari. Molto importante è anche l’effetto “spigolo”,
il quale si manifesta con una fragilità notevole della zona adiacente alla linea di incontro tra due superfici perpendicolari.
L’attività di ricerca qui presentata è stata svolta in collaborazione con Siderit s.r.l., che si è riproposta di ricercare soluzioni
alternative per la conduzione dei trattamenti di nitrurazione
gassosa, allo scopo di diminuire i tempi ciclo. Dopo un’analisi
accurata dei tradizionali processi produttivi, volta a caratteriz-
2
Altezza [cm]
400
200
Potenza [kW]
150
150
zare la loro efficacia nel garantire la ripetibilità nel tempo delle
proprietà conferite al materiale trattato, si è proceduto nello
studio della rielaborazione dei cicli di produzione al fine di
renderli più veloci nella loro realizzazione. I risultati ottenuti
con i cicli rielaborati sono stati accuratamente confrontati con i
“vecchi” cicli produttivi per verificarne, in particolare, la comparabilità metallurgica. Scopo ultimo è il miglioramento della
produttività, operando con cicli affidabili più brevi di quelli
standard, con la possibilità di eseguire un maggior numero di
trattamenti nel medesimo periodo di tempo. Infatti, alla sperimentazione e raccolta dei dati ottenuti ha fatto seguito anche
una valutazione di fattibilità e di riduzione dei tempi di processo per i cicli produttivi testati, sulla base di un confronto
qualitativo con i processi già applicati industrialmente.
MATERIALI E METODI
La scelta del materiale da utilizzare nell’indagine sperimentale
è stata effettuata a seguito di una attenta analisi dei materiali
più frequentemente trattati presso Siderit s.r.l.. È stato scelto
l’acciaio 42CrMo4, la cui composizione nominale è riportata
in Tab. 1. Tale acciaio, pur non essendo un tipico acciaio da
nitrurazione, è comunque in grado di diffondere azoto fino a
spessori efficaci elevati ed è ampiamente utilizzato per la costruzione di particolari meccanici soggetti a sollecitazioni cicliche e a sfregamenti (mozzi, perni, alberi, camme, ecc).
Sono stati presi in considerazione due cicli produttivi standard,
nominati NT3 ed NT4, relativi al trattamento di nitrurazione
gassosa che consente di ottenere uno spessore efficace rispettivamente di 0.3 mm e 0.4 mm, con tolleranza ±0.1 mm. La prima fase del lavoro ha riguardato la campionatura sistematica
di tali cicli produttivi, al fine di ricavarne la efficienza e la ripetibilità nel tempo in termini di durezza conferita alla superficie
del pezzo, adeguatezza dei gradienti di durezza nel substrato
materiale, quantità di coltre bianca e qualità della struttura
metallografica che deve risultare esente da reticoli continui di
nitruri nello strato di diffusione. I cicli relativi ai trattamenti
NT3 ed NT4 sono stati osservati per due impianti produttivi, le
cui caratteristiche principali sono riportate in Tab. 2.
Per entrambi i forno a pozzo sono stati ricavati i dati-ciclo dai
programmi memorizzati nei rispettivi controllori informatici
che li gestiscono in maniera automatica: temperature, tempi e
portate dei gas. Variando questi ultimi parametri è stato possibile ottenere i due cicli produttivi NT3 ed NT4, i cui diagrammi temperatura-tempo adottati nei due impianti sono riportati
in Fig. 1. I cicli di nitrurazione gassosa standard sono composti da due fasi: la prima, detta di “arricchimento”, si svolge
a 500°C, mentre la seconda, detta di “diffusione”, si svolge a
540°C. Nelle diverse fasi del trattamento sono stati controllati
sistematicamente i valori del grado di dissociazione dell’am-
Memorie >>
Trattamenti termici
s
Fig. 1
Diagrammi temperatura-tempo indicativi per i
cicli standard NT3 ed NT4.
Temperature-time diagrams for NT3 and NT4 standard
cycles.
moniaca. I cicli standard sono stati modificati ed i cicli sperimentali sono stati nominati rispettivamente NT3-N2O ed NT4N2O per metter in evidenza l’impiego del protossido d’azoto.
Le modifiche, rispetto ai cicli standard, hanno previsto un preriscaldamento del materiale a 380°C per uniformare il più possibile la temperatura dei pezzi, uno step di preossidazione del
materiale, allo scopo di attivare la superficie preventivamente
all’arricchimento, e l’inserimento del protossido d’azoto nella
successiva fase di arricchimento. Si è deciso di mantenere una
proporzione tra il protossido d’azoto e l’ammoniaca del 2-3%,
lasciando inalterata l’atmosfera nitrurante nella fase di diffusione. L’effetto della preossidazione dell’acciaio in atmosfere
nitruranti si esplica in un aumento della profondità di indurimento [14, 21]. La campionatura di ogni ciclo termochimico è
stata effettuata inserendo provini cilindrici del materiale sopra
descritto all’interno delle cariche di produzione. Le dimensioni dei provini sono di 20 mm di spessore e diametro di 35 mm,
ricavati dal taglio di una barra cilindrica dello stesso diametro trattata termicamente con processo di bonifica. La durezza
media del materiale base è di 310÷320 HV1. Prima di eseguire
la nitrurazione, i pezzi sono stati accuratamente puliti per eliminare ossidi ed eventuali lubrificanti che possono alterare le
condizioni di assorbimento di azoto sulla superficie; inoltre,
sulle facce da nitrurare ogni campione è stato preventivamente
trattato con operazioni di sgrossatura e successivamente finite
con lapidellatura accurata. Tutti i campioni, trattati secondo i
cicli standard e sperimentali con protossido d’azoto, sono stati
successivamente sottoposti a caratterizzazione metallografica
e a misure di durezza superficiale, di gradiente di durezza e
di spessore della coltre bianca. Il parametro di confronto più
importante che è stato preso in considerazione è lo “spessore
efficace”, definito dalla Norma UNI 5478.
CARATTERIZZAZIONE MECCANICA E
MICROSTRUTTURALE
I campioni, dopo essere stati estratti dai forni di trattamento,
sono stati accuratamente sezionati secondo le sezioni longitudinali e trasversali, e preparati per le indagini di laboratorio.
Tutte le successive analisi microstrutturali e meccaniche sono
s
Fig. 2
Rappresentazione grafica del gradiente di durezza HV1.
Diagram of HV1 hardness gradient.
state eseguite con strumentazioni dedicate e seguendo i principali riferimenti normativi [27-28].
Prove di microdurezza
Un’accurata indagine sclerometrica è stata eseguita sui campioni in modo da ricavare i dati utili alla caratterizzazione del
trattamento termochimico. Sulle sezioni trasversali sono state
effettuate numerose prove di durezza per definirne il gradiente. I risultati ottenuti sono stati raccolti in fogli di calcolo ed
espressi graficamente (Fig. 2). Tenendo in considerazione il
valore medio della durezza a nucleo, è possibile individuare
la profondità efficace ottenuta, definita dalla norme UNI 5478,
come quel valore di durezza superiore di 100 unità HV rispetto
alla durezza a nucleo.
L’indagine sclerometrica è stata estesa anche alla superficie
esterna direttamene nitrurata, con l’analisi della durezza superficiale. Per ogni campione sono state eseguite le prove di
durezza superficiale HV03, HV05, HV1, HV5, HV10, HV30.
L’esecuzione di prove di durezza superficiale con carico di
prova sempre maggiore è utilizzata per verificare a quale di
s
Fig. 3
Misura dello spessore medio della coltre bianca; attacco Nital2 (500X).
Measure of the average white layer thickness; Nital2
etching (500X).
3
Trattamenti termici
Processo
Programma
Cariche osservate
Parametri trattamento
Arricchimento (h)
Diffusione (h)
Totale (h)
Durezza superficiale HV1
Profondità efficace [mm]
Coltre bianca [μm]
Durezza a nucleo HV1
Impianto 1
NT4
NT3
24
6
16
2
Valori medi Valori medi
37
37
45
45
82
82
670÷690 670÷690
0.48÷0.50 0.3 ÷0.40
13÷15
10÷12
310÷320 310÷320
<< Memorie
Impianto 2
Impianto 1
Impianto 2
NT4
NT3
NT4-N2O NT3-N2O NT4-N2O NT3-N2O
15
6
29
11
29
11
5
4
8
1
4
5
Valori medi Valori medi Valori medi Valori medi Valori medi Valori medi
18
25
20
20
20
13
27
40
30
30
30
20
45
65
50
50
50
33
670÷690 670÷690 670÷690 670÷690 670÷690 670÷690
0.42÷0.45 0.38÷0.40 0.42÷0.45 0.35÷0.38 0.42÷0.45 0.35÷0.38
10÷12
10÷12
14÷15
14÷15
12÷14
12÷14
310÷320 310÷320 310÷320 310÷320 310÷320 310÷320
s
Tab. 3
Dati relativi al campionamento dei cicli standard e sperimentali.
Standard and innovative tests data.
questi carichi, e quindi indirettamente a quale pressione specifica, si abbia un cedimento per fragilità della coltre bianca, e
quindi una possibile scheggiatura. È stato riscontrato che l’influenza della coltre bianca sulla durezza risulta evidente per
carichi fino al 1kg.
Analisi metallografica
Oltre alle prove sclerometriche è stata effettuata una indagine
metallografica per valutare lo spessore della coltre bianca e la
microstruttura; si è agito eseguendo un attacco chimico sulla superficie accuratamente lucidata dei provini inglobati con
una soluzione di alcool puro e acido nitrico al 2% (NITAL 2).
Dalla Fig. 3 si nota la presenza di nitruri di forma lamellare, la
cui quantità è maggiore vicino al bordo esterno e diminuisce
verso la zona interna. Non sono state rilevate strutture a reticolo, particolarmente dannose per la resistenza meccanica del
pezzo nitrurato.
ANALISI DEI RISULTATI E DISCUSSIONE
I risultati di interesse, accuratamente elaborati e discussi, riguardano i valori di durezza superficiale, il valore della profondità efficace e la quantità di coltre bianca conferita dal
trattamento. I valori sono stati mediati sul numero di prove
eseguite, cercando di tenere conto dei molteplici fattori di influenza sul trattamento.
Campionatura cicli standard
Seguendo le metodologie descritte nei precedenti paragrafi,
sono stati campionati inizialmente i cicli produttivi già standardizzati dall’azienda. I cicli oggetto di analisi sono stati
quelli aventi come obiettivo il raggiungimento di un valore
di profondità efficace di indurimento di 0.3 mm (NT3) e di
0.4 mm (NT4). Questi cicli vengono differenziati in funzione
dell’impianto osservato.
Campionatura cicli con protossido d’azoto
Durante l’esecuzione della campionatura dei cicli standard
si sono affrontate anche le prime sperimentazioni relative al
protossido di azoto. La sperimentazione è stata soggetta ad
evoluzioni per quanto riguarda la conformazione dei cicli sperimentali applicati, in considerazione del fatto che il materia-
4
le trattato faceva parte della produzione aziendale ordinaria.
Spesso sono stati confrontati i dati relativi ai campioni della
sperimentazione con analisi eseguite sui pezzi meccanici commerciali trattati, al fine di verificare l’efficacia dei cicli anche
su di essi.
Tutti i dati, di cui la Tab. 3 è un parziale consuntivo, sono stati
raccolti ed elaborati con fogli elettronici per poter eseguire una
approfondita indagine di tipo statistico.
Confronto tra le caratteristiche fisico-meccaniche
Durezze superficiali
L’analisi statistica ha messo in evidenza che i valori medi della
durezza superficiale, valutata con carichi crescenti, sono perfettamente confrontabili tra i cicli standard e quelli sperimentali
omologhi con protossido; le differenze di durezza superficiale
ottenuta per i cicli omologhi sono entro campi di tolleranza del
tutto accettabili. Nei grafici di Fig. 4 si pone in evidenza anche
il confronto tra i due impianti. Le prove di durezza superficiale
sui campioni sottoposti a nitrurazione con protossido di azoto non hanno generato criccatura evidente della coltre bianca
nemmeno con carichi elevati. Minime criccature sono state rilevate solo nei rari casi di coltre bianca con spessori oltre i 17
μm.
Gradienti di durezza
Complessivamente si rileva che i cicli con protossido di azoto
hanno portato a buoni risultati di durezza efficace. Di seguito,
nelle figure 5-6, sono riportati i gradienti di durezza medi ottenuti per i trattamenti NT3 ed NT4 con e senza protossido di
azoto, nei due impianti oggetto di sperimentazione; i gradienti
sono associati ai programmi di ciclo di trattamento, codificati
nel controllore informatico.
In Fig. 5a sono riportati i gradienti medi di due cicli sperimentali di trattamento, effettuati con l’impianto 1; si può notare
che il ciclo standard (P.24) garantisce una profondità efficace
media di 0.5mm, il gradiente medio relativo al primo ciclo sperimentale (P.29) arriva a 0.45mm efficaci e il secondo ciclo sperimentale (P.28) consente di arrivare a 0.5mm efficaci. Durante
la sperimentazione, l’osservazione del risultato medio ottenibile con i cicli standard ha messo in evidenza la possibilità di
ridurre la profondità efficace per non conferire ai pezzi un eccessivo indurimento. Pertanto, la produzione è stata indirizza-
Memorie >>
Trattamenti termici
a
a
b
b
s
Fig. 4
s
Fig. 5
Grafico di confronto della durezza superficiale
ottenuta negli impianti 1 e 2: (a) per i cicli NT4 ed NT4N2O, (b) per i cicli NT3 ed NT3-N2O.
Surface hardness comparison obtained in implants 1 and 2:
(a) NT4 and NT4-N2O cycles, (b) NT3 and NT3-N2O cycles.
Confronto tra i gradienti di durezza medi
per cicli NT4 ed NT4-N2O: (a) impianto 1, (b)
impianto 2.
Average hardness gradient comparison for NT4 and
NT4-N2O cycles: (a) implant 1, (b) implant 2.
ta all’uso del programma (P.29) che garantisce una profondità
efficace di 0.45 mm. Di queste considerazioni si terrà conto nel
confronto tra i tempi ciclo, nel paragrafo dedicato a questo tipo
di analisi.
Per l’impianto 2 i risultati ottenuti con i cicli standard NT4 e
con quelli sperimentali NT4-N2O sono perfettamente comparabili (Fig. 5b). Il ciclo sperimentale (P.29) fornisce profondità
efficaci di 0.45 mm comparabili con i cicli standard (P.15).
I cicli sperimentali per la classe di trattamento NT3 hanno
anch’essi dato buoni risultati nel raggiungimento della profondità efficace, anche se la profondità efficace raggiunta dai
cicli standard è leggermente più elevata per entrambi gli impianti (Fig. 6). Questi risultati possono essere legati al fatto che
la diminuzione percentuale di tempo apportata ai cicli NT3
standard, per entrambi gli impianti, è leggermente superiore
rispetto alla diminuzione apportata ai cicli NT4. Inoltre, la sperimentazione relativa ai cicli di trattamento tipo NT3 è stata
più limitata per via del fatto che i due impianti oggetto della
sperimentazione stessa sono prevalentemente impiegati per
l’esecuzione di trattamenti di nitrurazione profonda NT4. È da
osservare, infine, che per gli stessi cicli standard NT3 le profondità raggiunte sono risultate essere addirittura eccessive,
quindi la leggera differenza di profondità può essere considerata del tutto tollerabile se non desiderabile.
È stato osservato che i trattamenti sperimentali hanno garantito una buona ripetibilità dei gradienti di durezza. Dal punto
di vista produttivo è di estrema importanza la capacità di fornire un risultato ripetibile e quindi controllabile. In Fig. 7 si
riportano dei grafici che mettono in evidenza la ripetibilità dei
gradienti di durezza per il ciclo sperimentale NT4-N2O, rispettivamente nell’impianto 1 e nell’impianto 2 .
Si osserva che i gradienti di durezza sono dispersi entro un
campo di ampiezza del tutto accettabile (4-5%). Si può quindi
affermare che la sperimentazione abbia confermato una buona
ripetibilità dei gradienti di durezza ottenibili con i cicli di nitrurazione con protossido.
Strato dei composti
L’impiego del protossido di azoto, accelerando la cinetica di
nitrurazione, ha una influenza sulla formazione e sulla crescita
degli strati induriti. Si è verificato che un uso inadeguato dello
stesso protossido porta a spessori della coltre bianca troppo
elevati, e quindi indesiderabili, pertanto deve essere ben controllato per non portare a risultati finali non voluti. Per i due
diversi impianti, si riporta di seguito la statistica relativa agli
spessori della coltre bianca e dello strato poroso; sono messi in evidenza i valori medi ottenuti con i vari programmi di
nitrurazione, e vengono confrontati con i cicli omologhi che
5
Trattamenti termici
<< Memorie
a
a
b
b
s
Fig. 6
s
Fig. 7
Confronto tra i gradienti di durezza medi
per cicli NT3 ed NT3-N2O: (a) impianto 1, (b)
impianto 2.
Average hardness gradient comparison for NT3 and
NT3-N2O cycles: (a) implant 1, (b) implant 2.
Ripetibilità dei gradienti di durezza per i
trattamenti sperimentali NT4-N2O: (a) impianto 1,
(b) impianto 2.
Repeatability of hardness gradients for NT4-N2O
innovative treatments: (a) implant 1, (b) implant 2.
prevedono l’impiego di protossido di azoto.
In Tab. 4 si può notare che i valori medi degli spessori delle
coltri bianche ottenute con i cicli che prevedono l’impiego di
protossido sono entro limiti più che accettabili, seppure leggermente superiori, ma del tutto confrontabili, con i rispettivi
cicli omologhi che non impiegano protossido. In ogni caso gli
spessori di coltre bianca rilevati non sono mai risultati superiori ai limiti di accettabilità imposti dalla norma UNI 5478. Nei
grafici di Fig. 8 è riportata la dispersione dei dati relativi agli
spessori delle coltri bianche e degli strati porosi rilevati nella
sperimentazione per due cicli di nitrurazione NT4 ed NT4N2O eseguiti nell’impianto 1.
TRATTAMENTO
Impianto
Programma
Coltre bianca [μm]
Strato poroso [μm]
NT4
1
24
13.5
7
NT4 -N2O
1
29
14.5
7.5
NT4
2
15
12.5
6
Durata dei trattamenti – Confronto tra cicli omologhi
La sperimentazione ha evidenziato la possibilità di ridurre i
tempi-ciclo per i trattamenti analizzati. Le progressive piccole
variazioni dei cicli hanno permesso di ottimizzare i tempi-ciclo
in funzione del risultato richiesto dalla produzione aziendale.
Il confronto dei tempi-ciclo deve essere effettuato tenendo in
considerazione il risultato finale ottenuto. In Fig. 9 si riportano
NT4 -N2O
2
29
13
7.5
NT3
1
6
11
4.5
NT3 -N2O
1
11
13.5
7.5
NT3
2
6
10
5
NT3 -N2O
2
11
12.5
6.5
s
Tab. 4
Riassunto dei dati relativi agli spessori di coltre bianca ottenuta per i cicli di trattamento NT4 ed NT4-N2O,
NT3 ed NT3-N2O.
Data referred to the white layer thicknesses in NT4, NT4-N2O, NT3 and NT3-N2O treatment cycles.
6
Trattamenti termici
Memorie >>
TRATTAMENTO
Programma
Impianto
Tempo ciclo [h]
Profondita’ efficace [mm]
Risparmio tempo
NT4
24
1
82,5
0,5
NT4 -N2O
28
1
62
0,5
NT4
15
2
65
0,45
25%
NT4 -N2O
29
2
52
0,45
20%
NT3
6
1
45
0,4
NT3 -N2O
11
1
32
0,38
NT3
6
2
45
0,4
29%
NT3 -N2O
11
1
32
0,38
27%
s
Tab. 5
b
a
s
Riassunto dei tempi-ciclo e calcolo del risparmio percentuale a parità di risultato per i trattamenti NT4, NT4-N2O, NT3 ed
NT3-N2O per gli impianti 1 e 2.
Time-cycle data and percentage saving evaluation in NT4, NT4-N2O, NT3 and NT3-N2O treatments for implants 1 and 2, in the same
operational conditions.
Fig. 8
Spessori
della coltre bianca
e del relativo strato
poroso per i cicli di
nitrurazione a) NT4
e b) NT4-N2O in
impianto 1.
Thicknesses of the
white layer and of the
corresponding nitrides
substrate for nitriding
cycles a) NT4 and b)
NT4-N2O in implant 1.
a confronto i tempi-ciclo per i programmi (P.24) relativo al trattamento NT4, e i programmi (P.28) e (P.29) per il trattamento
NT4-N2O; si può notare che il ciclo derivato dal programma
(P.29) riduce drasticamente il tempo di svolgimento del trattamento, facendo risparmiare 30 h; tuttavia la profondità efficace
ottenibile con tale ciclo è inferiore rispetto al ciclo standard, da
quanto rilevato in 4.3.2. Si riporta in Tab. 5 un confronto tra i
tempi-ciclo, posti in relazione a parità di risultato ottenuto.
Per i cicli NT4-N2O il risparmio di tempo-ciclo è evidente e
può essere considerato attorno al 20÷25%, rispetto ai cicli standard, tenendo dovutamente conto degli errori sperimentali.
Per quanto riguarda i cicli NT3-N2O anche in questo caso si riscontra un marcato risparmio di tempo, addirittura superiore
rispetto ai cicli più lunghi. Questo potrebbe essere dato dal fatto che l’adsorbimento di azoto, e quindi la nitrurazione, non ha
un andamento lineare con il procedere del trattamento; infatti,
l’ispessimento della coltre bianca e l’allontanamento del fronte
di adsorbimento di azoto rendono più difficoltoso il passaggio
dell’azoto stesso. Dunque la capacità di adsorbimento di azoto
cala con lo svilupparsi del trattamento, e i cicli più corti risentono maggiormente del favorevole impiego del protossido di
azoto. Da non sottovalutare è inoltre la possibile diversa risposta del forno all’impiego del protossido di azoto; la diversa
capacità dissociativa è dipendente anche dalla geometria della
storta, nonché dal moto di ricircolo al suo interno. È possibile che le diversità tra gli impianti portino a diverse efficienze
s
Fig. 9
Grafici temperatura-tempo di confronto per
alcuni programmi di nitrurazione NT4 e NT4-N2O.
Temperature-time diagrams: comparison between NT4
and NT4-N2O nitriding programs.
7
Trattamenti termici
<< Memorie
a
Costo [€/kg]*
Peso specifico [kg/l]
NH3
0.496
0.00076
N2O
3.35
0.00198
s
Tab. 6
Parametri di riferimento per il calcolo dei costi (*dati
relativi al 2006).
Reference parameters for cost evaluation (*2006 data).
b
s
Fig. 10
Confronto tra i tempi-ciclo standard e quelli
sperimentali: a) impianto 1, b) impianto 2.
Comparison between standard and innovative time-cycle:
a) implant 1, b) implant 2.
dei cicli sperimentali. Nei grafici di Fig. 10 si pone in evidenza
il risparmio sui tempi-ciclo, ottenuto con la sperimentazione,
per i singoli impianti.
VALUTAZIONI ECONOMICHE
Analisi dei costi per i cicli produttivi
Questa analisi si basa essenzialmente sul confronto tra i costi
di processo dei cicli per i trattamenti standard e per quelli che
prevedono l’uso del protossido d’azoto, prendendo in considerazione, per quanto possibile, i fattori maggiormente incidenti.
Vi sono fattori energetici, relativi al dispendio/contenimento
dei gas tecnici di processo, e fattori legati alla pianificazione
della produzione.
Costi dei gas di processo
La stima dei costi dei gas di processo nei cicli di trattamento è
stata condotta calcolando la quantità di gas tecnici impiegati,
almeno teoricamente, nei trattamenti omologhi. Sono stati presi in considerazione i parametri relativi ai costi e ai pesi specifici dei gas tecnici riportati in Tab. 6.
È stato sviluppato un algoritmo di calcolo che in funzione dei
tempi e sulle portate dei gas di processo è in grado di valutare il
risparmio di un ciclo rispetto all’altro. Analizzando con tale algoritmo il costo del ciclo NT4 standard secondo il programma
(P.24) e del ciclo sperimentale NT4-N2O secondo il programma (P.28), che hanno dato in fase di sperimentazione risultati
8
operativi analoghi, è stato stimato un risparmio percentuale
pari al 18% per il trattamento sperimentale. Nonostante, infatti, il maggiore costo specifico del protossido di azoto rispetto
all’ammoniaca, il ciclo sperimentale ha un minore costo complessivo rispetto a quello standard grazie all’abbattimento dei
tempi-ciclo.
Costi dell’energia elettrica
La riduzione dei tempi-ciclo porta anche ad un minor dispendio in termini di energia elettrica spesa. Sono stati presi in considerazione i costi energetici delle varie fasi di riscaldamento e
di mantenimento per i cicli standard e per quelli sperimentali
con protossido.
Ponendo a confronto i costi complessivi di energia elettrica,
valutati per il programma (P.24) del ciclo standard NT4 e per il
programma (P.28) del ciclo sperimentale NT4-N2O, si è stimato
un risparmio energetico pari a circa il 16%. È da osservare che
il risparmio sui costi dovuti all’energia elettrica è comparabile
con quello ottenuto sul costo dei gas tecnici.
CONCLUSIONI
Nel presente lavoro, che ha costituito l’attività sperimentale di
una tesi di Laurea in Ingegneria dei Materiali, è stato eseguito
un confronto tra cicli standard di trattamento termochimico di
nitrurazione gassosa e cicli sperimentali con protossido d’azoto, eseguiti sull’acciaio 42CrMo4. Dal confronto tra i cicli standard e quelli sperimentali è stato messo in evidenza come le
caratteristiche tecniche dei campioni trattati e analizzati siano
risultate del tutto comparabili per ciò che riguarda la durezza
superficiale, gli spessori di coltre bianca ottenuti e la morfologia dei nitruri come rilevato delle analisi microstrutturali. Il
vantaggio più evidente riguarda il risparmio di tempo-ciclo a
parità di profondità efficace ottenuta, confermato ampiamente
dall’utilizzo sistematico dei cicli testati anche nella produzione effettiva. Ulteriori vantaggi economici risultano evidenti
dal confronto dei costi di produzione. Il maggior costo specifico del protossido di azoto è ammortizzato ampiamente dal
risparmio complessivo sui gas grazie alla diminuzione dei
tempi-ciclo e quindi ad un minore impiego complessivo di
ammoniaca. Inoltre, il calo dei tempi-ciclo porta ad una non
trascurabile diminuzione dei consumi di energia elettrica per
effetto di una conseguente minore necessità di mantenimento
alle elevate temperature di processo. Infine, la diminuzione dei
tempi di impegno impianto ne apporta una maggiore disponibilità, traducibile in una possibile migliore capacità di pianificazione della produzione.
RINGRAZIAMENTI
Si ringrazia l’azienda Siderit s.r.l di Zola Predosa (BO), tutto il
personale operativo presente nel laboratorio metallurgico e nei
reparti produttivi, in particolare il responsabile del progetto di
Memorie >>
Trattamenti termici
ricerca Sig. Paolo Porcu, per aver reso possibili le sperimentazioni oggetto di questo lavoro.
BIBLIOGRAFIA
1] D. GHIGLIONE et al., “Pratique des traitements thermochimiques – Nitruration, nitrocarburation et dérivés”, Techniques
de l’ingénieur, traité Matériaux Métalliques M 1 doc.227 (1996)
2] P.SCHAAF, “Laser nitriding of metals”, Progress in material
science 47 (2002) p.1-161
3] E. LEHRER, “Über das Eisen Wasserstoff-Ammoniak“ ,
Gleichgewicht Z. Elektrochemie 36-6 (1930) p.382-392
4] A. BURDESE, “Metallurgia e tecnologia dei materiali metallici”, UTET, Torino (1992), p.424
5] T. BELL, ASM Handbook vol.4 , ed. ASM, Metals Park , Ohio
(1991), p.387
6] C. CIBALDI, “I criteri di scelta e di trattamento degli acciai
da costruzione e da utensili”, AQM, Brescia (2006), p.503
7] J.D. FAST, “Interactions of metals and gases”, Kinetics and
Mechanisms 2, Macmillan Press. LTD., London (1971)
8] M. JKORWIN and A. MANCUSO, “La nitrurazione controllata in gas”, Rivista di meccanica (II) 9 (1987) p.180
9] W. LERCHE, Freiberger Forschungshefte 185 (1976) p.1
10] M. ROSSO et al., “Innovative nitriding treatment applied
to PM steels”, Proc. 11th Congr. of the Int. Fed. for Heat Treat.
and Surf. Eng. and 4th ASM Heat Treat. and Surf. Eng. Conf. in
Europe, ed. AIM, Milano, (1998), p.403-412
11] B. PRENOSIL, Härterei Technische Mitteilungen 20 (1965)
p.41-49
12] H. J. ECKSTEIN and W. LERCHE, Neue Hütte 14 (1968)
p.210-215
13] J. SLYCKE and L. SPROGE, “Kinetics of the Gaseous Nitrocarburising Process”, Surface Engineering 5-2 (1989) p.125140
14] F. CAVALLERI et al., La Metallurgia Italiana 82 (1990)
p.599-604
15] S. FOISSEY et al., Proc. “11th congr. of the IFHT and 4th
ASM conf. in Europe, Firenze, 19-21 0ctober (1998), ed. AIM,
Milano, p.291-300
16] F.M. MONTEVECCHI, “Evoluzione dei criteri di regolazione e controllo dei processi di nitrurazione degli acciai”, La
Metallurgia Italiana 82-6 (1990) p.605-614
17] P. SCHAAF and F. LANDRY, “Mössbauer Investigation of
Nitriding Process - Gas Nitriding and Laser Nitriding”, MSMS
(1998)
18] J. KAZIOR et al., “Thermochemical Treatment of Fe–Cr–
Mo alloys”, Surface and Coatings Technology 151-152 (2002)
p.333-337
19] H. J. GRABKE et al., “Nitridation in NH3-H2O Mixtures” ,
Materials and Corrosion 54-11 (2003) p.895-901
20] H. J. GRABKE, “A case of Nitridation, Carburization and
Oxidation on a Stainless Steel”, Materials and Corrosion 55-6
(2005) p.384-388
21] A. TIKHONOV, “Methods of Acceleration of Saturation
Processes During Carbonitriding Treatment”, La Metallurgia
Italiana 4 (2005) p.33-37
22] H.J. SPIES et al., “Influence of Oxygen Additions during
Gas Nitriding on the Structure of the Nitride Layers”, Materialwissenschaft und Werkstofftechnik Iss. 10-29 (1998) p.588594
23] A. BRAMLEY et al., “The Diffusion of Non-Metallic Elements in Iron and Steel”, Transactions of the Faraday Society
5-31 (Maggio 1935) p.707-734
24] P. HAYES and P. GRIEVESON, “The effect of phosphorus
and oxygen on the nitriding of α-Fe”, Acta Metallurgica 23 8
(1975)
25] J.P. PEYRE et al., “Nitruration par le procédé Alnat N, caractéristiques du procédé, structure des couches réalisées”,
Traitement Thermique n° 219, (Maggio 1988)
26] J.P. PEYRE et al., “Caractéristiques mécaniques des couches nitrurées par le procédé Alnat N”, Traitement Thermique
227 (Aprile 1989)
27] Norma UNI 5478, Trattamenti termici dei materiali metallici – Nitrurazione - Terza edizione, (Settembre 1999)
28] Norma UNI EN ISO 6507/1, Prove di durezza Vickers Metodo di prova, (Luglio 1999)
ABSTRACT
USE OF NITROGEN PROTOXIDE IN THE
THERMOCHEMICAL GAS NITRIDING TREATMENT: STUDY
OF PRODUCTION PROCESSES AND METALLURGICAL
CHARACTERISATION
Keywords: nitrogen protoxide, surface hardening,
metallography, gas nitriding
The gas nitriding treatment is usually carried out in an anhydrous ammonia environment inside retort furnaces at 500-550°C. The process times are
variable according to the desirable depth of surface hardening, but in general are very time consuming. This scenario motivated this research work,
in particular studying the possibility to increase the industrial process rate.
Many mechanical components in structural steels are subjected to this treatment in order to obtain an high surface hardening. The surface hardening
is due to the deposition of ε iron nitrides directly on the metal surface and
to the precipitation of γ iron nitrides on the substrate. The research for new
efficient and industrially applicable solutions led to use an additional gas in
the process. In the present research the effect of the introduction of nitrogen
protoxide within the anhydrous ammonia in the process environment has
been tested. Tests have been performed inside of the Siderit s.r.l. establishments in Zola Predosa (Bologna-Italy). Initially, the standard nitriding
processes in the anhydrous ammonia environment have been studied, introducing samples in 42CrMo4 steel within the production materials inside
of the retort furnaces. The cycles able to guarantee surface hardening of
0.3 and 0.4 millimeter depths - namely NT3 and NT4 respectively - have
been taken into account. Subsequently, the same samples have been realised
introducing the protoxide of nitrogen as a processing gas and evaluating
its quantity respect to ammonia, step by step. The mechanical and metallurgical characterisations of the standard processes and of the ones with
the use of N2O within the nitriding environment have been performed; Vickers hardness tests on the samples’ treated surfaces and also Vickers micro
hardness profiles on the treated thicknesses have been carried out. Moreover, the quality of the nitriding layers, obtained by the different processing
parameters, have been evaluated through a careful metallographic analysis.
The use of nitrogen protoxide can improve the kinetic of the process, which
leads to a reduction of the process times respect to the standard processes
of about 20-30%. Despite that, the mechanical and metallurgical characteristics of the samples treated with the additional N2O gas are comparable
with the ones treated by means of the traditional thermochemical process.
This way, the use of nitrogen protoxide could be of great importance in
industrial nitriding processes. Pros and contras in terms of feasibility and
cost are discussed on the basis of qualitative comparisons with currently
applied industrial processes.
9
Trattamenti termici
Memorie >>
INFLUENCE OF HEAT TREATMENT
ON THE MICROSTRUCTURE AND
TOUGHNESS OF BÖHLER M333
ISOPLAST STEEL
J. Perko, C. Redl, H. Leitner
In this work the through hardenability and the influence of the heat treatment parameters (austenitizing
temperature, cooling parameter λ and tempering temperature) on the microstructure and the achievable
toughness level of Böhler M333 ISOPLAST are investigated. The results are compared to the standardized
tool steel grade DIN 1.2083. The investigations showed that the cooling parameter λ has a strong influence
on the impact toughness of M333 ISOPLAST plastic mould steel. The toughness is reduced by pro-eutectoid
precipitates and not by a lack of through hardenability. Furthermore, it was found out that depending on the
cross section of the moulds appropriate heat treatments lead to a good combination of hardness and toughness.
KEYWORDS: plastic mould steel, heat treatment, toughness, through hardenability, pro-eutectoid precipitates
INTRODUCTION
MATERIALS AND INVESTIGATION METHODS
Plastic mould steels which are currently available on the market, e.g. DIN 1.2083, 1.4028 and 1.2316, are often not able to
fulfil the high requirements of the plastics processing industry. For that reason, Böhler Edelstahl developed the pressureelectro-slag remelted, nitrogen alloyed tool steel grade M333
ISOPLAST. Nitrogen has a lot of positive effects on martensitic
chromium steels [1, 2]. The partial replacement of carbon by
nitrogen leads to an increase in corrosion resistance and toughness. General corrosion is reduced as well as pitting and crevice corrosion. The improvement in toughness results primarily
from the very homogeneous distribution of fine precipitates
[3]. Thus, M333 ISOPLAST combines excellent mirror finish
polishability with highest cleanliness and toughness levels
and excellent corrosion properties.
However, as a consequence of the growing demand for large
moulds, plastic mould steels must also exhibit an excellent
through hardenability in order to avoid the formation of bainite
or pearlite during quenching. Additionally, due to lower cooling rates a low tendency to form grain boundary precipitates,
which cause grain boundary embrittlement, is required.
Therefore, this work concentrates on the investigation of the
through hardenability and on the influence of the heat treatment parameters on the mechanical properties and microstructure of M333 ISOPLAST.
Tab. 1 shows the chemical composition of M333 in comparison
to M310, which approximates the standardized tool steel grade
DIN 1.2083. This steel was used as a reference steel grade for
the dilatometer investigations. The samples for the dilatometer
investigations of M333 and M310 were manufactured from a
hot-rolled and soft-annealed bar with a diameter of 86 mm. For
the dilatometer experiments a quenching dilatometer Bähr Dil
805 A/D was used. To follow the evolution of the hardness of
the dilatometer samples, Vickers hardness values (HV10) were
measured using microhardness tester supplied by Zwick.
For the samples for impact toughness testing of M333 two
slices with a thickness of 60 mm were cut from the top of a
forged and soft annealed bar with the dimension 603 x 303
mm2. Then, the samples were cut in longitudinal direction at
half radius and were heat treated with an oversize of 0.5 mm
on every side. For the hardening of the samples a vacuum heat
treatment furnace was used. Quenching was performed by using nitrogen as quenching gas.
The cooling parameter λ is defined as the cooling time from
800 to 500 °C in seconds divided by 100. The selected cooling
parameters λ were adjusted with dummy samples exhibiting
the same size as the test specimens. The temperature was controlled by mounting a thermocouple in a drilled hole in the
centre of the dummy sample. The heat-treated samples were
grinded to the final dimension of 7 x 10 x 55 mm3. The impact
tests were performed with a 450 J pendulum Roell Amsler 101.
Four samples of each heat treatment were tested and the average and standard deviation were calculated. Rockwell C hardness was measured on each specimen using a hardness tester
Emco-Test M4R 025 G3. The fracture surface of all tested im-
J. Perko, C. Redl
Böhler Edelstahl GmbH & Co KG, Kapfenberg, Austria
H. Leitner
University of Leoben, Austria
1
Trattamenti termici
Steel grade
M310 ISOPLAST
M333 ISOPLAST
DIN
~1.2083
-
C
0.38
0.28
Si
0.70
0.30
Mn
0.30
0.30
<< Memorie
Cr
14.30
13.50
V
0.20
-
N
+
s
Tab. 1
Chemical composition of the tool steels investigated.
Composizione chimica degli acciai da utensile investigati.
Austenitizing temperature [°C]
Holding time [min]
Cooling parameter λ [s x 10-2]
Tempering temperature [°C]
Tempering time [min]
980/1020
30
0.5/5/8
250/300/510/520/ 530/540
2 x 120
s
Tab. 2
Applied heat treatments for impact toughness
testing.
Trattamenti termici applicati per la verifica della resilienza.
s
Fig. 1
RESULTS AND DISCUSSION
The aim of the dilatometer investigations was to find out how the
austenitizing temperature and the cooling parameter λ influence
the transformation behaviour of M333 in comparison to M310.
During heating and cooling of a sample in the dilatometer, the
change in length (ΔL) as a function of time and temperature is recorded, from which ΔL/L0 versus temperature (T) is plotted. L0
represents the original length of the sample. Representative plots
of ΔL/L0 versus T of M310 and M333 at the two austenitizing
temperatures selected (980 °C and 1020 °C) are shown in Fig. 1.
It can be seen that in case of M333 for both austenitizing temperatures and all applied cooling rates only one transformation
takes place during cooling. This is indicated by the distinct length
change starting at about 250 °C. Microstructural investigations
revealed that this length change is caused by the transformation
austenite to martensite. However, the martensite start temperature (MS) is shifted to lower temperatures with increasing austenitizing temperature, which is due to the higher solution state of
the austenite at higher temperatures. When M333 is austenitized
at 980 °C, the remaining change in length after quenching is lower
M333 ISOPLAST
Cooling parameter λ
980 C
1020 C
HV10
MS [°C]
HV10
MS [°C]
0.5
593
300
636
278
8
567
342
628
285
30
537
370
601
279
* BS … Bainite Start Temperature
2
Dilatometer curves of M310 and M333.
Curve dilatometriche degli acciai M310 e M333.
compared to 1020 °C. This indicates that the amount of retained
austenite is less compared to austenitizing temperature 1020 °C,
what was confirmed by XRD measurements.
In contrast, the steel grade M310 shows different transformation
behaviour. At an austenitizing temperature of 980 °C and the
highest investigated cooling parameter λ=30 additional transformations can be observed, starting at about 750 °C and 400 °C.
Microstructural examinations of this sample showed the presence
of pearlite and bainite. Thus, an austenitizing temperature of 980
°C is not suitable for parts with large cross sections made of this
steel grade. For an austenitizing temperature of 1020 °C, M310
shows neither a pearlitic nor a bainitic transformation at a cooling parameter λ=30. However, for both steel grades the measured
MS-temperatures are summarized in Tab. 3.
The experimentally determined hardness measurements of the
dilatometer samples are also listed in Tab. 3. In case of M333 Tab.
3 reveals that while the hardness decreases with increasing λ
values, the MS-temperature increases. This indicates a decreasing solid solution state with increasing λ values which might be
due to the formation of pro-eutectoid phases such as carbides or
carbonitrides.
The microstructure of selected dilatometer samples is shown
in Fig. 2. In order to reveal the microstructure, Nital was used
as etchand. In Fig. 2, the former austenite grain boundaries can
M310 ISOPLAST
0.5
8
30
563
509
295
317
340
406*
617
582
503
281
335
369
s
pact specimens was macroscopically documented. Additionally, the fracture surface of selected heat treatment cycles was
analyzed using a scanning electron microscope Jeol JSM6460lv
(SEM). Tab. 2 shows an overview of the applied heat treatments of the samples for impact toughness testing.
Tab. 3
MS-temperature
and hardness of M333
and M310 after various
heat treatment cycles.
Temperatura MS e durezza
degli acciai M333 e M310
dopo differenti cicli di
trattamento termico.
Memorie >>
Trattamenti termici
s
Fig. 3
s
Fig. 2
Impact energy and hardness of M333, austenitized
at 980 °C.
Valori di resilienza e durezza dell’acciaio M333, austenizzato a
980 °C.
Microstructure of selected heat treatment cycles.
Microstruttuta relativa ai diversi cicli di trattamento termico
scelti.
not be seen in both M333 samples quenched with λ=0.5. On the
other hand, the grain boundaries of M310 at an austenitizing temperature of 1020 °C are weakly visible, indicating that the grain
boundaries are decorated. Additionally, M310 contains some
coarser M23C6 carbides, which is one of the reasons for the significantly lower toughness level in comparison to M333, as reported in [4, 5].
At a cooling parameter λ=8 both steels show a strong decoration of
the grain boundaries, which is designated by the distinct appearance of the grain boundaries. However, the strongest decoration
shows M333 when it is austenitized at 1020 °C. The M333 samples
austenitized at 980 °C exhibit a finer grain and the decoration of
the grain boundaries is less pronounced. The reason for the lower
austenite grain size of the M333 samples austenitized at 980 °C
is the higher content of undissolved precipitates, which stabilize
the grain size during austenitizing. At 1020 °C their amount is
strongly reduced, which causes grain growth.
The reason for the decoration of the former austenite grain
boundaries seems to be due to the formation of pro-eutectoid,
chromium-rich precipitates as reported in [6, 7]. Herzog [7] investigated oil and air quenched samples with SANS and found
out that the amount of pro-eutectoid precipitates in the size range
from 1 to 100 nm increases with increasing austenitizing temperature and cooling parameter, which correlates well with the results
of M333.
The results of the dilatometer investigations indicated, that the interaction of pro-eutectoid precipitations and austenite grain size
may have more impact on the toughness of M333 than a lack of
through hardenability, as the grain boundary decoration occurred
at λ-values below 8. Therefore, the impact toughness investigation of M333 concentrated on λ-values below 8.
However, the investigations are primarily focused on the heat
treatment of larger cross sections. For that reason the tempering
of most samples was carried out at temperatures between 510
and 540 °C, which are higher than the secondary hardening peak
of M333. Temperatures above the secondary hardening peak are
needed to reduce residual stresses and retained austenite as far as
s
Fig. 4
Impact energy and hardness of M333, austenitized
at 1020 °C.
Valori di resilienza e durezza dell’acciaio M333, austenizzato a
1020 °C.
possible. However, for comparison the tempering temperatures
250 and 300 °C were included in the investigations.
Fig. 3 and 4 show the impact energy versus tempering temperature of the samples austenitized at 980 °C and 1020 °C, respectively. Additionally, the measured hardness values are plotted in the
diagrams. The results reveal that the hardness at each tempering
temperature is virtually independent of the cooling parameters
up to λ=8. The hardness level of the samples hardened at 980 °C is
continuous 1-2 HRc lower than the hardness of samples austenitized at 1020° C. Nevertheless, a hardness of approximately 50
HRc, which is typical for the plastic moulding industry, is still
achievable with an austenitizing temperature of 980 °C.
In contrast to the hardness, the achievable toughness level is
strongly depending on the cooling parameter λ, especially for the
samples austenitized at 1020 °C. As long as the cooling parameter λ is low (λ=0.5), the samples austenitized at 1020 °C show
high hardness and toughness values in all heat treatment cycles
3
Trattamenti termici
s
Fig. 5
Impact energy of M333 as a function of the cooling
parameter λ.
Valori di resilienza dell’acciaio M333 in funzione del parametro
di raffreddamento λ.
investigated. When the cooling parameter is increased to λ=5, the
toughness is significantly reduced, especially for the most relevant tempering temperatures 510 and 520 °C, which have to be
applied to reach hardness levels of 50 HRc at temperatures above
the secondary hardening peak.
For the most relevant tempering temperatures for industrial application, the impact energy as a function of the cooling parameter λ is illustrated in Fig. 5. It can be seen that the cooling parameter λ has only little influence on the impact toughness of samples
austenitized at 980 °C, which is beneficial for the heat treatment
of tools with large cross sections. In contrast to that, the negative influence of the cooling parameter λ on the impact toughness
of samples austenitized at 1020 °C is more pronounced. For that
reason, an austenitizing temperature of 1020 °C seems to be appropriate for tools with small cross sections only.
To investigate the microstructural mechanisms that cause embrittlement of M333, the fracture surface of the tested specimens was
investigated macroscopically and by means of SEM (Fig. 6 and
<< Memorie
7).
Fig. 6 shows that the cooling parameter λ has only little influence on the fracture appearance of samples austenitized at 980°C.
There is only a marginal reduction of ductility when samples
with low and high cooling parameters are compared, which is
in accordance with the measured impact toughness values. In
contrast to that, the samples austenitized at 1020°C show a high
loss of ductility with increasing λ values. This effect is extremely
pronounced at tempering temperatures of 510 and 520 °C, where
the fracture mode changes from ductile to mainly brittle already
at λ=5. This is also reflected in the low impact toughness values
measured.
A detailed fracture analysis was conducted by SEM investigations. Fig. 7 shows the fracture surface of the samples austenitized at 980 and 1020°C and tempered at 510 °C in the region of
the incipient crack.
It is obvious that both samples quenched with λ=0.5 exhibit a
ductile fracture, which corresponds well with the high impact
toughness values measured. When the cooling parameter λ is increased up to 5 the fracture appearance changes significantly. The
samples which were austenitized at 980°C still exhibit a ductile
fracture, but the intergranular character increases. In contrast to
that, the samples austenitized at 1020 °C and cooled with λ=5 exhibit an entire intergranular fracture with no ductile character at
all, being the reason for the low impact toughness of these samples. When the cooling parameter is increased to λ=8, the intergranular character of samples austenitized at 980 °C increments,
but there is still a high portion of ductile fracture visible.
A more detailed view on the surface of the fracture of the samples hardened at 1020°C with cooling parameter λ=5 confirms
this theory (see Fig.8). However, the presence of pro-eutectoid
precipitates could not be clearly confirmed with SEM-EDX, as
the resolution of this method is too low. But it can be assumed
that the formation of precipitates preferentially takes place at the
grain boundaries, as the critical nucleation energy is lower there
and because of the lower surface energy that needs to be brought
up. Therefore, precipitations with a low coherency develop preferentially there [6].
However, the investigations revealed that a lack of through hardenability can be excluded; the embrittlement with increasing λ
values can only be caused by a disadvantageous combination
of austenite grain size and high solution state of the austenite.
This encourages the formation of pro-eutectoid precipitates at the
austenite grain boundaries at higher λ values, which causes the
s
Fig. 6
s
Fig. 7
Macroscopic comparison of the broken M333
samples.
Confronto macroscopico della superficie di rottura del provino in
acciaio M333.
SEM images of the fracture surfaces (all samples
were tempered at 510 °C).
Immagini al SEM delle superfici di rottura (tutti i campioni sono
stati rinvenuti a 510°C).
4
Trattamenti termici
Memorie >>
s
Fig. 8
Detailed view of the SEM image in Fig. 7 (1020
°C / λ=5).
Particolare dell’immagine al SEM di Fig. 7 (1020 °C / λ=5).
observed embrittlement of M333, especially in the region of the
secondary hardening peak.
Additionally, heat treatment simulations were carried out in order to determine the temperature in the centre of a block during quenching. Nitrogen gas with a pressure of 5 bar was used
as quenching medium. Fig. 9 depicts the temperature evolution
in the centre of work pieces with different square cross sections.
The results show that a cooling parameter λ of approximately 8 is
achievable in the centre of a block with a dimension between 150
x 150 mm2 and 200 x 200 mm2. However, Fig. 9 also reveals that
cooling of real work pieces does not follow one λ-value. At the beginning cooling is always slower, which enables the formation of
pro-eutectoid precipitates on grain boundaries. This cooling characteristic also applies for the quenching of the impact toughness
samples investigated. Therefore, the measured impact toughness
values reflect the behaviour of real work pieces.
CONCLUSIONS
The cooling parameter λ has a strong influence on the impact
toughness of M333 ISOPLAST plastic mould steel. It was found
that the toughness is reduced by pro-eutectoid precipitates and
not by a lack of through hardenability.
If low cooling parameters λ can be attained, e.g. during heat
treatment of tools with small cross sections, austenitizing at 1020
°C results in the best combination of high hardness and impact
toughness. On the other hand, the application of lower austenitizing temperatures, e.g. 980 °C, significantly reduces the negative
influence of higher cooling parameters λ on the impact toughness.
This study showed that Böhler M333 ISOPLAST is a suitable steel
grade for moulds with small and large cross sections, if an appropriate heat treatment is applied. However, additional work is
required to gain detailed information on formation of the precipi-
s
Fig. 9
Temperature evolution in the block centre for
different dimensions (N2 5 bar).
Evoluzione della temperatura al centro del blocco per diverse
dimensioni (N2 5 bar).
tates which cause embrittlement.
ACKNOWLEDGMENTS
The authors are grateful to Carlos Martinez, Böhler Edelstahl
GmbH & Co KG, for performing the SEM studies.
REFERENCES
1] R. SCHNEIDER, F. KOCH, P. WÜRZINGER, Proc. of the 5th International Tooling Conference, Leoben, September 1999, p.265.
2] G. LICHTENEGGER, Proceedings of the 15th Leobener Kunststoffkolloquium, Leoben, November 1999, p.XIV-9.
3] G. LICHTENEGGER, R. SCHNEIDER, J. SAMMER, G. SCHIRNINGER, P. WÜRZINGER, J. NEUHERZ, Proc. of the 5th International Tooling Conference, Leoben, September 1999, p.644.
4] K. SAMMT, J. SAMMER, J. GECKLE, W. LIEBFAHRT, Proc.
of the 5th International Tooling Conference, Karlstad University,
September 2002, p.285.
5] D. UGUES, M. ROSSO, A. RIVOLTA, Proc. of the 6th International Tooling Conference, Politecnico di Torino, May 2006,
p.325.
6] H. BERNS, M. POHL, Praktische Metallographie 18 (1981),
p.209.
7] C. HERZOG, Master thesis, University of Leoben (1995).
ABSTRACT
INFLUENZA DEL TRATTAMENTO TERMICO SU
MICROSTRUTTURA E TENACITÀ DELL’ ACCIAIO
BOHLER M333 ISOPLAST
Parole chiave: trattamenti termici, proprietà, acciaio
In questo lavoro sono stati analizzati temprabilità e l’influenza dei parametri di trattamento termico (temperatura di austenizzazione, parametro
di raffreddamento λ, temperatura di rinvenimento) sulla microstruttura
e sul livello di durezza ottenibile per l’acciaio Böhler M333 ISOPLAST.
I risultati sono stati confrontati con il classico acciaio per utenisili DIN
1.2083. Le indagini hanno evidenziato che il parametro di raffreddamento λ ha una forte influenza sulla resistenza all’urto dell’ acciaio M333
ISOPLAST per stampi destinati all’industria delle materie plastiche. La
tenacità risulta ridotta dai precipitati pro-eutettoidici e non dalla mancanza di temprabilità. Inoltre, si è accertato che a seconda della sezione
trasversale degli stampi, i trattamenti termici appropriati portano ad un
buon compromesso fra durezza e tenacità.
5
Magnesio e leghe
Memorie >>
NEW CHALLENGES AND DIRECTIONS
FOR HIGH PRESSURE DIE-CAST
MAGNESIUM
L. Zaffaina, R. Alain, F. Bonollo, Z. Fan
Being the lightest structural metallic material, Magnesium is widely used by the automotive industry when
searching for solutions to address energy consumption, emission reductions and vehicles handling. The need to
use light-weight solutions is becoming stronger than it has been so far, as other technologies allowing to reduce
emissions have already yield a lot of their potential.
This wider use of Magnesium requires to grow more knowledge on the material properties, corrosion protection
techniques and improved production processes. However material R&D cannot be limited to property studies
and potentials: directions for innovation implies to understand all the benefits of the material and process for
each application and their values to the customer.
In this work, the components’ key attributes will be highlighted, together with the efforts to lower costs and the
potential use of new technologies.
KEYWORDS: magnesium, optimisation, applications, automotive, CAE led design, innovation, value analysis, HPDC,
melt conditioned HPDC
INTRODUCTION
Over the last decades, the demand for magnesium components
in the automotive industry kept increasing. The key drivers
have been the need to reduce fuel consumption and emissions
and the need to improve vehicle handling. However there
are other advantages to use cast magnesium solution such as
improved tolerances, higher potential to contribute to vehicle
stiffness, higher fit and finish, improved noise vibration and
harshness performance; but these are difficult to evaluate and
not often quantified in assessing the overall product values.
In order to plot the directions for innovation in terms of technology and applications, it would be useful to clearly understand the list of benefits of magnesium casting and their value
to the customer.
A view of the market from Meridian’s point of view is shown
here, together with the technologies whose have triggered new
applications to appear in the market.
applications are transfer cases and transmission housings
with more than 2 million cast per year; steering column
brackets and instrumentation panels are the main body
component produced. (respectively 3 and 2,5 million parts
per year).
New components are gaining growing interest at car makers, as indicate increased number of request for quotations
and ongoing development studies; theses applications are
Front End Structures, Engine Cradles, Centre Consoles and
A MARKET VIEW
Fig. 1 shows the main cast magnesium components for automotive industry. Meridian’s key products in power-train
Luca Zaffaina
Università di Padova, Italy
Regis Alain
Meridian Technologies Inc, Italy
Franco Bonollo
Università di Padova, Italy
Zhongyun Fan
Brunel University, United Kingdom
s
Fig. 1
Main automotive diecast magnesium applications.
Principali componenti in magnesio prossocolato nell’auto
la metallurgia italiana >> marzo 2009
25
Magnesio e leghe
<< Memorie
Power train applications (transmission cases, cam covers,
oil pan, engine brackets).
The general message perceived from the costumers is that
the need to use light-weight solutions is becoming stronger than it has been so far, as other technologies allowing
to reduce emissions have already yield a lot of their potential.
these are not always easy to capture by a tier 1 or tier 2
supplier. For a similar product, the functions to achieve can
vary a lot depending on the customers or vehicle and the
benefit of using Magnesium will vary. Each product been
specific even within a family of products, general rules on
reason to use Magnesium are difficult to extract.
More focus has been put to try to understand drivers to
chose Magnesium, what influences costs and where the
costumer will find more interest in the technology. In other
words where will the costumer find values. According to
J.C. Anderson and J. A. Naurus “Value is the worth in monetary terms of the technical, economic, service and social
benefits a customer company receives in exchange for the
price it pays for a marketing offer” [3].
It is important to keep in mind that the customer does not
really want materials or services. He wants Use and/or
Aesthetic functions to be accomplished. [4]
On products two kind of functions can be individuated:
pre-requisites and attributes.
The pre-requisites are the functions for which there is a target to reach and the customer is not willing to see value in
an extra-performance.
At the opposite, for attributes, the customer is willing to
value an increased performance. By extension the costumer
could also review his requirements at the system level and
allow a reduction in performance in exchange for a price
reduction.
CHANGES ON THE MARKET AND THEIR EFFECT
WORK TO LOWER THE COST OF FUNCTIONS
With regards to magnesium applications, the recent changes that have happened are as follow:
a) Availability of chromium-free corrosion prevention solutions helped to introduce exterior components with aesthetic requirements like Front End Carriers, door frames,
roof mechanisms or rear window frames.
b) Low creep alloys with high ductility (AE44) led to the
introduction of magnesium castings for engine applications, where both crash and high temperatures have to be
met. Meridian produces the first example of engine cradle
for the General Motors Corvette Z06 and more components
are currently under development. In addition to the high
temperature mechanical properties, AE44 shows better corrosion resistance than common used ductile alloys (AM60,
AM50): the general corrosion resistance is even higher than
the AZ91’s commonly used for transfer cases (see Fig. 2).
As this part is not cosmetic, just like a transfer case, the Engine Cradle is not coated. [1]
c) Over the last two years, the cost of Aluminium alloys
increased while magnesium alloys remained stable, hence
Magnesium has become more attractive for power train
components. This cost evolution and the increased need for
weight saving is driving new interest and developments
from costumers.
As all automotive suppliers, Meridian works to reduce
costs, and different approach are used. The usual one,
based on product cost breakdown, is not sufficient to identify actions; it has been found that more ways to reduce cost
could be yield by doing the analysis operating on the cost
per function. An extension of this approach is the identification of other markets i.e. new applications for which current technology could reach other values at an interesting
cost.
In an ideal world, it would be useful to plot for each function, the state of the market, i.e. for each function, what level
is reached, at which costs, where Magnesium technologies
are present. A successive step should be to assess if different values could be reached, at which costs and with which
technology improvements. Unfortunately such approach is
not feasible as it would require to grab a high amount of
information from the customer.
For a single component, the system is complex, and it is actually sometime not even obvious for the customer to know
that he has the information about component/system functions and their value: in fact the data may lay in different
functional areas [3].
Considering the complexity of the component and the system, a simple approach has to be developed so that it can be
reapplied to other products.
The investigation has then to focus on:
- identify functions
- understand attributes and values
- assess the cost for attributes and values
- reduce cost to achieve attributes and values
- understand costumers sensitivity to different values
In this work, the component’s key attributes will be highlighted, together with the efforts to lower costs and the potential use of new technologies.
s
Fig. 2
Various Mg and Al alloys performance after
240hours ASTM B117 salt spray [2].
Confronto tra le prestazioni di alcune leghe di Mg e Al
dopo 240 ore in nebbia salina (ASTM B117)[2].
UNDERSTANDING WHAT CUSTOMERS VALUE
The changes mentioned above are not sufficient to understand the growth of specific applications on the market
and why for some costumers Magnesium has become core
roots for defined applications, while others do not follow
the trend.
Above weight saving, costumers see other interesting features in the use of Magnesium on specific applications:
26
marzo 2009 << la metallurgia italiana
Magnesio e leghe
Memorie >>
s
Fig. 3
s
Fig. 4
Sample of AM60B after 1000h of salt spray test ASTM B117.
Provino in AM60B dopo 1000 ore di nebbia salina - ASTM
B117.
Evolution of FEC weight and performance.
Evoluzione del peso e delle prestazioni dei FEC.
NEW COMPONENTS
a) Front End Carrier (FEC)
The features that costumers value for front ends are:
- weight saving (40 to 50%), i.e. 3 to 9 kg depending on the type
of vehicle [5]
- induced weight saving: the reduced weight at the front of the
vehicle reduces load on the body in white
- better fit and finish performance compared to steel when requirements increase
- high stiffness achievable, hence potential to redistribute stiffness target at the system level
- better air flow management offered by foundry shape
- higher integration offered by the foundry compared to welded
technologies
In the past, the biggest issue facing the use of magnesium for
exterior applications has been corrosion prevention and in particular galvanic corrosion. Today galvanic and cosmetic corrosion requirements are achieved by proper design of component
and fixings and proper choice of joint materials. Two principles
are to respect. The first one is to insure a good drainage in fixing
areas to avoid that mud and humidity get trapped. The second is to separate Magnesium from steel by the introduction of
an Aluminium spacer, respecting a simple design rule that is to
have a minimum distance of 10 mm between Magnesium and
steel.
Cosmetic targets have to be achieved, hence parts have to be
coated. Current solution consist of an acid etch to remove surface iron particles picked up during casting process, a chemical
conversion coating to insure the adhesion of the final top coat
that is an epoxy base powder coating. This solution is capable of
1000 hours salt spray tests and of 12 weeks humidity salt spray
tests (see Fig. 3).
To reduce cost, new development consider top plastic cover so
that the magnesium part is not visible anymore and coating can
be removed.
Over last six years, the use of magnesium for FEC extended
from light truck to SUV and finally to passenger car where component weight decreased from 5 to 3 kg, whilst increasing crash
s
Fig. 5
FECs developments and production trend.
Nuovi FEC in sviluppo e in produzione.
requirements were met. (Fig. 4)
The potential of Front End Structure is strongly dependant from
the type of Front End and the functions it has to achieve.
The number of part in production in Europe has increased from
1 to 4 from 2004 to 2007 and 7 will be in production in 2009.
Components under development and physical validation increase coincidently, as shown in Fig. 5.
The first development was carried out in North America and led
to development and production of the Corvette Z06 Engine Cradle (Fig. 6). The project’s aim was to achieve significant weight
savings and it led to the development of a new alloy capable
of high to moderate strength, ductile behaviour to function in
cyclic loading, elevated temperature and corrosive environment
resistance.
The magnesium solution led to a 32% weight saving in comparison with aluminium which means 4,8kg. The new alloy is capable of working at 125°C, has improved elongation (9-10% with
7mm-thick walls) and good corrosion resistance. Proper alloy
properties together with an accurate fixings design allowed to
avoid any surface treatment. For this sport car the key attribute
27
Magnesio e leghe
s
Fig. 6
Corvette Z06 Engine Cradle.
Supporto motore della Corvette Z06.
<< Memorie
is an important weight saving associated with the size of the
part in an area where Magnesium had not been used before.
Alloys development has allowed new chassis and power
train application to be considered: productions processes
in conjunction with material properties and Meridian design approach [6] will lead to more engine cradles, engine
brackets , power train components.
c) Centre Console Structure (CCS)
The Centre Console Structure is one of these products for
which requirement can differ strongly for one application
to another. Key attributes are:
- weight saving of 50% compared to steel
- high integration in a small mass of Magnesium, hence
very cost efficient
- increased potential to offer storage volume offered by
casting potential
- improved assembly precision due to one-piece solutions,
which improves visible perception of quality
- can be structural and transfer load in lateral impact in specific vehicle configurations
- can be cosmetic with a coating: an example of that is the
ford GT40 console
Costumers show more and more interest in Centre Console
Structure, and the amount of development is growing significantly.
INCOMING APPLICATIONS AND TECHNOLOGIES
s
Fig. 7
Microstructure of AZ91 alloy bar (6mm diameter) produced by the MC-HPDC process showing fine
and uniform microstructures and close to zero porosity
(<0.2%).
Microstruttura di una barra di AZ91 (diametro 6mm)
prodotta con il processo MC-HPDC. Si notano la microstruttura fine a uniforme e la porosità prossima a zero
(<0,2%).
s
Fig. 8
Comparison of mechanical properties of AZ91D
alloy with different levels of scrap addition produced by
the MC-HPDC process.
Confronto delle caratteristiche meccaniche della lega
AZ91D prodotta con il processo MC-HPDC utilizzando
differenti percentuali di materiale di scarto.
28
Some interesting opportunities and projects under development are briefly mentioned here.
a) Increased use of optimisers
CAE-led design approach [7] gives interesting advantages
such a significant reduction of time to deliver the concept,
possibility to provide to customer the CAE model in advance to verify structural performance and a mid-surface
model to check the package.
In addition to that, the use of optimisers since the early
concept-design stage makes that geometry can be design
to maximise the value to customer by using optimisers to
lower the cost for functions or to minimise the weight-saving premium price.
Ongoing and future developments will extend this approach with increased use of optimisers on process and
design.
b) MC-HPDC process
Melt conditioned high pressure diecasting (MC-HPDC)
process, previously named rheo-diecasting (RDC) process,
is a combination of a MCAST (melt conditioning by advanced shear Technology) machine for melt conditioning
with a standard high pressure diecasting (HPDC) machine
for component shaping [8].
Laboratory trials show that the MC-HPDC process is capable of producing close to zero porosity samples with a fine
and uniform microstructure throughout the entire component (Fig. 7). MC-HPDC process provides samples with improved strength and ductility in the as-cast condition, compared with HPDC or other available semisolid processing
techniques. Initial tensile test have shown elongations up to
20% on AM50 test sample [9].
Extended work on this process has shown that the MCAST
machine is actually mostly acting as a grain refining device
and that it could be used with 100% liquid, meaning that
all process parameters can remain unchanged compared
to conventional HPDC process. The addition of this single
marzo 2009 << la metallurgia italiana
Magnesio e leghe
Memorie >>
piece equipment to an existing HPDC machine would limit
investment and permits new applications. Potential applications could be in the suspension area, together with AE44
showing high temperature resistance, very good corrosion
performance and increased ductility.
Another interesting characteristic of the MCAST machine is
the fact that it breaks up oxides in small round-shaped particles well dispersed through the entire components. This
characteristic opens another potential that is the possibility
to use scrap parts and runners systems for specific internal
applications, leading to interesting cost savings [10]. Fig.
8 compares the mechanical properties of AZ91 alloy produced by the MC-HPDC process using different amount of
scraps.
c) New trend on approaching cosmetic corrosion
As market grows, cost for coating has significantly decreased; another consideration is the fact that the coating is
acting just to respect visual requirements, while the structural-galvanic corrosion resistance is obtained through
spacers and proper design of components. Hence in some
cases where part would not be visible coating could be removed: it could be the case on FEC by masking the top
surface with a plastic cover.
REFERENCES
[1] N. LI, R. OSBORNE, B. COX, D. PERNOD “Magnesium Engine Cradle – The USCAR structural Cast Magnesium Devel-
opment Project” SAE Technical Paper Series 2005-01-0337
[2] Hydro Brochure “Corrosion Properties and Protection”
05-2005.
[3] J.C. ANDERSON, J.A. NAURUS “Business marketing:
understand what customers value” Harvard Business Review, November-December 1998
[4] L.D. MILES “Principle of Value Analysis”, 1976
[5] A. TIPPINGS, T. LAWSON, “Design and development
of a High Pressure Die “FEC” (Front End Carrier) For Exterior Automotive Applications” 62ND Annual World Magnesium Conference, 2005
[6] R. ALAIN, T. LAWSON et al. “Robustness of large thin
wall magnesium die castings for crash applications” SAE
Technical Paper Series 2004-01-0131
[7] T. LAWSON “Magnesium applications in modern vehicle applications” Euromotor Conference “New Advances
in Body Engineering” 2007
[8] Z. FAN, “Development of the Rheo-diecasting process
for Mg-Alloys” Material Science and Engineering A 413414 (2005) pag. 72-78
[9] S. JI, Z. ZHEN and Z. FAN “Effects of rheo-die casting
process on the microstructure and mechanical properties of
AM50 magnesium alloy” Material Science and Technology,
21(2005),1019-1024
[10] G. LIU, Y. WANG and Z FAN: A physical approach
to the direct recycling of Mg-alloy scrap by the rheo-diecasting process, Materials Science and Engineering: A, 472
(2008), 251-257
ABSTRACT
NUOVE SFIDE E DIREZIONI DI SVILUPPO PER IL
MAGNESIO PRESSOCOLATO
Parole chiave: magnesio e leghe, corrosione, pressocolata, tixoforatura, impieghi alta/bassa temperatura,
processi, produzione, proprietà
È noto come il magnesio sia il più leggero tra i materiali strutturali:
è questo il principale motivo per cui è utilizzato in maniera estensiva
nell’industria automobilistica (fig. 1) al fine di minimizzare i consumi
energetici e le emissioni inquinanti e nel caso in cui si voglia migliorare
la maneggevolezza dei veicoli.
La spinta verso l’alleggerimento ha visto un’accelerazione negli ultimi
anni dovuta probabilmente al fatto che altre tecnologie tese a ridurre le
emissioni hanno già raggiunto gran parte del loro potenziale.
Un utilizzo più ampio del Magnesio comporta la necessità di aumentare
la conoscenza delle proprietà del materiale, delle tecniche di protezione
dalla corrosione (figg. 2, 3) e di processi produttivi migliorati.
Tuttavia le attività di R&S non possono essere limitate allo studio delle proprietà e dei potenziali del materiale: indirizzare l’innovazione in
termini di tecnologie e applicazioni implica comprendere tutti i benefici
del materiale e del processo per ogni applicazione e il loro valore per il
cliente.
Oltre alla riduzione di peso ci sono infatti altri vantaggi nell’uso di soluzioni in magnesio pressocolato quali: riduzione delle tolleranze, potenzialità di contribuire alla rigidezza del veicolo, miglioramento dell’assorbimento delle vibrazioni; purtroppo questi benefici sono difficili da
valutare e spesso non quantificati nella stima del valore complessivo del
prodotto da parte del cliente.
In questo lavoro sono analizzati alcuni dei cambiamenti tecnologici e di
mercato che hanno permesso l’estensione del magnesio pressocolato a
nuove applicazioni (fig. 6) e il loro impatto su specifici componenti.
Si è evidenziato come i cambiamenti citati non siano sufficienti a spiegare la crescita di specifiche applicazioni (fig. 5) e come sia difficile estrarre
regole generali che spieghino le ragioni per cui si scelga di utilizzare o
meno soluzioni in magnesio pressocolato.
Diventa quindi importante lavorare al fine di comprendere quali siano le
funzioni del componente e i loro costi e cercare di indirizzare le attività
di R&S alla riduzione del costo delle specifiche funzioni.
A titolo esemplificativo sono sottolineati in questo lavoro gli attributi
chiave per alcuni componenti (fig. 4), gli sforzi tesi a ridurne il costo e il
potenziale d’uso di alcune nuove tecnologie (figg. 7, 8).
29

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