F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
465–473
NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF
ADDITIVE MANUFACTURING VERSUS CONVENTIONAL LPS
POWDER METALLURGY
PROIZVODNJA KERMETOV NA OSNOVI NBC; PRIMERJAVA MED
DODAJALNO TEHNOLOGIJO L-PBF IN KONVENCIONALNO
METALURGIJO PRAHOV LPS
Fabio Miranda
1,3*
, Marcelo Otavio dos Santos
1,2
, Daniel Rodrigues
3
,
Rodrigo Santiago Coelho
4
, Gilmar Ferreira Batalha
1
1
University of São Paulo, Polytechnic School, EPUSP–PMR, São Paulo, Brazil
2
Mauá Institute of Technology, University Centre, São Caetano do Sul, Brazil
3
BRATS Sintered Filters and Metallic Powders, Cajamar-SP, Brazil
4
SENAI CIMATEC – Institute of Innovation for Forming & Joining of Materials, Salvador, Bahia, Brazil
Prejem rokopisa – received: 2023-01-30; sprejem za objavo – accepted for publication: 2023-08-28
doi:10.17222/mit.2023.972
The production of carbide parts (cermet) by additive manufacturing, such as laser powder bed fusion (L-PBF), has been a great
challenge due to the complex optimization of process parameters to improve density, porosity, microcracks or abnormal growth
of grains and obtain a microstructure as homogeneous as possible. This work aims to compare the evolution of the
microstructure when using the conventional route of powder metallurgy, i.e., liquid phase sintering (LPS) with the L-PBF direct
additive manufacturing process, considering the NbC-based carbide material. Sample compositions were prepared in w/%, sam-
ples were compacted under 50–125 MPa, without polymeric binders, and they were sintered under a vacuum at temperatures of
1330 °C and 1370 °C. For the L-PBF process, a vibrating device made it possible to improve the fluidity of the mixtures of three
alloys, NbC–30Co, NbC–30Ni and NbC–30(Co, Ni). The mixtures exhibited low sphericity, low fluidity and compressibility,
which were improved with a roller compactor. Thin powder mixture deposition layers were evenly applied and well distributed
across the powder bed to avoid defects and cracks during sintering. The L-PBF process parameters varied including a power of
50–125 W and a laser scanning speed of 25–125 mm·s
–1
. Different microstructures, identified with a light microscope (LM) and
a scanning electron microscope (SEM), and properties obtained with the two processes, direct (L–PBF) and indirect sintering
(LPS), were compared.
Keywords: NbC–cemented carbides, L–PBF additive manufacturing, liquid phase sintering, microstructure and mechanical
properties
Proizvodnja karbidnih trdin oziroma kompozitov na osnovi drobnih karbidnih delcev v kovinski osnovi (kermeti; kompoziti
keramika-kovina) z dodajalno tehnologijo, kot je naprimer lasersko pretaljevanje nanosov plasti prahu za plastjo (L-PBF; angl.:
Laser Powder Bed Fusion), je zelo zahtevno zaradi komplicirane optimizacije procesnih parametrov, ki naj izbolj{ajo gostoto,
zmanj{ajo poroznost in {tevilo mikrorazpok ali nenormalno (pretirano) rast kristalnih zrn in s tem zagotovijo ~imbolj homogeno
mikrostrukturo. V tem ~lanku avtorji opisujejo primerjavo razvoja mikrostrukture karbidne trdine na osnovi NbC nastale s
postopkom konvencionalne metalurgije prahov oz. postopkom sintranja v prisotnosti teko~e faze (LPS; angl.: Liquid Phase
Sintering) z mikrostrukturo nastalo med direktnim L-PBF postopkom. Avtorji so surove vzorce za LPS pripravili s stiskanjem
pri 50 MPa do 125 MPa brez dodatka polimernega veziva in jih sintrali v vakuumu pri temperaturah 1330 °C in 1370 °C. Za
L-PBF postopek so uporabili vibracijsko napravo, ki je izbolj{ala teko~nost (nasipanje) treh izbranih pra{nih me{anic:
NbC-30Co, NbC-30Ni in NbC-30(Co, Ni). Te me{anice imajo delce dokaj neprevilne oblike in se zato slabo nasipavajo plast za
plastjo. Izbolj{anje zgo{~evanja plasti so dosegli z dodatno uporabo valja (angl.: roller compactor). Tako so bile enakomerno
nane{ene tanke plasti pra{nih me{anic z dobro porazdelitvijo pra{nih delcev v posteljici, kar naj bi prepre~ilo nastajanje razpok
in drugih napak med sintranjem. Parametre L-PBF procesa so varirali; mo~ med 50 W in 125 W ter skeniranje z laserjem s
hitrostjo od 25 mm·s
–1
do 125 mm·s
–1
Nato so avtorji izvedli {e primerjavo mikrostruktur in morfolo{kih lastnosti dobljenih z
obema postopkoma (direktnim L-PBF in indirektnim LPS postopkom). Primerjave so izvedli s pomo~jo svetlobnega oziroma
opti~nega (LM) in vrsti~nega elektronskega mikroskopa (SEM).
Klju~ne besede: karbidne trdine na osnovi NbC, dodajalna tehnologija (L-PBF), sintranje v teko~i fazi (LPS), mikrostruktura in
mehanske lastnosti
1 INTRODUCTION
New cermets based on NbC with alternative Ni
binder phases and a Co/Ni combination, being face-cen-
tered cubic (FCC), emerged to minimize the consump-
tion of traditional WC–Co alloys, mainly for the manu-
facture of cutting tools for machining subject to high
temperatures and wear resistance.
1
NbC
y
(0.77<y<0.96)
has been used as a secondary hard phase for high-speed
steels and as an inhibitor for the WC grain growth for
cermet tools, but its potential to be used as a primary
hard phase needs to be further explored.
3
Little attention
was given to NbC in the 19
th
century, but recent studies
have shown promising results, allowing us to replace
WC with NbC, especially for cutting tools and wear re-
sistance.
1–4
An NbC-based cermet with an amount of
Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473 465
UDK 546.261:536.421.5:52-334.2 ISSN 1580-2949
Original scientific article/Izvirni znanstveni ~lanek MTAEC9, 57(5)465(2023)
*Corresponding author's e-mail:
fabio.miranda@usp.br (Fabio Miranda)

combined carbon being above 10 w/%, includes NbC +
12Ni + C (free), with its eutectic temperature close to
1265 °C, as seen on Figures 1a and 1b
4
. The liquid phase
of FCC (Ni) starts at this temperature for amounts of
combined carbon of 9.33 w/% and 10 w/%. Below
9.25 w/%, there is slow cooling; temperatures of 1500 °C
and 1100 °C result in equilibrium phases, NbC–liquid
and NbC–NbNi
3
, respectively. With an amount of com-
bined carbon of 9.25–9.75 w/%, when being cooled
down from 1250 °C, the formation of an NbNi
3
com-
pound occurs, and there is a pseudo-binary region NbC +
FCC + NbNi
3
, with a very narrow range.
3,4
NbC
y
can
form sub-cubic carbides, allowing a wide range of
stoichiometry of the combined carbon, from
6.5–11.8 w/%, and can be adapted by stoichiometry as
shown in the Nb-C binary phase diagram, Figure 1d.
5
In terms of additive manufacturing, L–PBF is a
promising alternative for refractory alloys where metal
powders are selectively melted layer-by-layer using a
computer-controlled laser beam.
6
One advantage of the
L–PBF process is the ability to manufacture components
of complex shapes and small dimensions, something that
cannot be easily manufactured with the conventional pro-
cess of powder metallurgy. This technology has been
used more frequently by several industrial sectors in or-
der to reduce production steps, mainly for manufacturing
mechanical components with complex geometries, and
with a high value and low volume of production.
7
The
present work aims to compare the microstructure charac-
teristics and properties of NbC–30Co, NbC–30Ni and
NbC–30(Co, Ni) alloys, manufactured using the conven-
tional liquid phase sintering (LPS) process and laser
powder bed fusion (L–PBF), to point out the challenges
in the additive manufacturing of the cermet based on
NbC with alternative binding phases.
2 EXPERIMENTAL PART
This section presents the experimental procedures in-
cluding the conventional (pressing and sintering) and se-
lective laser melting AM (L–PBF) process. The follow-
ing cemented carbide samples were prepared, with the
compositions shown in Table 2: (I) NbC–30Co, (II)
NbC–30Ni and (III) NbC–30(Co, Ni). High-purity com-
F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
466 Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473
Figure 1: a) Phase evolution with temperature of the studied NbC-12Ni cermet, b) Phase diagram simulation of the NbC-12Ni system, c) Isother-
mal section of Nb-C-Ni at 1420 °C
4
, d) Binary Nb-C phase diagram
5

mercial Co, Ni and WC powders were used, with the fol-
lowing particle sizes: (2.0, 4.8 and 1.5) μm, respectively.
The metal powders of NbC, Co, Ni and WC were mixed
and homogenized by means of a conventional paddle
stirrer for 2 h, using isopropyl alcohol to avoid breaking,
crushing or reducing the particulates. Figure 2 shows the
SE of each mixture (Table 2) of the constituent powders
(Table 1) used in this work. The LPS samples were com-
pacted, without polymeric lubricants, into a metallic ma-
trix with a rectangular shape of (6 × 12 × 37) mm, using
a pressure of 50–125 MPa. Green compacted samples
were sintered in a vacuum oven (2 × 10
–2
mbar), with
sintering cycles at 1330 °C (1603 K) and 1370 °C
(1643 K) and a heating rate of 10 °C min
–1
(283 K/min),
then stabilized for 60 min and slowly cooled inside the
oven.
The L–PBF process was carried out in an Omni-
Sint-160 machine, as shown in Figure 3a. In this ma-
chine, the powder bed is created by a vibrating device
that has a rubber scraper ruler and a metallic roller. The
rotational speed of the roller is independently controlled
and should be zero. The area of the powder bed is too
large (around 630 cm
2
), and the compaction force is ap-
plied bya2k groller (19.6 N); thus, the compaction
pressure is very low. This powder feeding device was de-
veloped specially for this L–PBF machine. The metallic
F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473 467
Table 1: Properties of refractory metallic and ceramic powders
Powders
Theoretical
density
Average grain
size
Purity
Apparent
density Origin
(g·cm
–3
) (μm) (w/%) (g cm
–3
)
NbC 7.750 < 1.0 99.650 3.0 F&X Electro-Materials Limited (China)
WC 15.630 1.5 99.530 4.0 Buffalo Tungsten Inc. (USA)
Co 8.908 2.0 99.970 2.0 Nanjing Hanrui Cobalt Co. (China)
Ni 8.900 5.0 99.850 2.5 CVMR Corporation (Canada)
Table 2: Weight balance (w/%) of the prepared alloys
Alloy Description of cermets w/% Theoretical density
(g·cm
–3
) WC NbC Co Ni
I NbC–30Co 3.0 67.0 30.0 0.0 8.003
II NbC–30Ni 3.0 67.0 0.0 30.0 8.005
III NbC–30(Co, Ni) 3.0 67.0 15.0 15.0 8.004
Figure 3: a) OmniSint-160 machine, b) vibration device with a roller, c) sintering chamber, d) powder bed flowability & compressibility,
e) sintering with different parameter levels P
L
–v
s
Figure 2: SE images of the mixtures of metal powders: a) (I) NbC–30Co, b) (II) NbC–30Ni and c) III NbC–30(Co,Ni) alloy

powder is fed manually into a container that has a capac-
ity of 1500 g and is vibrating. It makes the powder go
through a 35 mesh sieve. The translation and rotation
movement simultaneously spreads and compacts the
powder, generating a powder bed, Figure 3d. Thin layers
are produced with the rollers and it is possible there is
even extra fine metal powder, used for 3D printing addi-
tive manufacturing. During direct sintering, the chamber
was inertized with a high-purity dynamic argon atmo-
sphere (0.3 L min
–1
), Figure 3c. The oxygen concentra-
tion during the process was around 300 μg/g. The depo-
sition was limited to 40 layers, with a layer thickness of
30 μm. The shrinkage during consolidating was high, so
the final sample thickness was smaller than 1200 μm.
Figure 4a shows a schematic of the process parame-
ters for L–PBF, Figure 4b presents a 3D view for the
L–PBF process, and Figure 4c presents the scanning
strategy with a 67° rotation, from layer N to layer N-1.
The laser power (P
L
) ranged from 50–125 W, the scan
speed (vs) parameter ranged from 50–150 mm·s
–1
, keep-
ing the scan line spacing (hs) equal to 80 μm. The layer
F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
468 Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473
Figure 4: a) Pressing method
8
, b) Zigzag scanning strategy, c) Monolayer deposition strategy with 67° rotation, d) Sintering strategy via L–PBF,
generated by different levels of factors P
L
–v
s
Figure 5: Types of apparent porosities (ISO 4505) and microstructures (ISO 4499) of samples NbC–30Co, NbC–30Ni and NbC–30(Co, Ni) pro-
duced via the conventional LPS route ( = 125 MPa)

thickness (Iz) was 30 μm. The focus diameter was
140 μm.
3 RESULTS AND DISCUSSIONS
The apparent porosities (LM 100×) and microstruc-
tures of the samples obtained with light microscopy
(LM) after the LPS process at temperatures of 1330 °C
(1603 K) and 1370 °C (1643 K) are presented in Fig-
ure 5. During sintering, solid-solid bonding occurs be-
tween the NbC carbides, before the emergence of the liq-
uid phase, forming necks between the particles and
porosities. The Ni and Co–Ni binding phases of the sam-
ples sintered at 1370 °C indicate good wettability around
the NbC grains. For the Co-based alloys, the NbC grains
are not well rounded. In the NbC–Co alloy, density can
be achieved in the solid-solid and solid-liquid state; how-
ever, the microstructures are not spheroidized and the
mechanical properties are very low.
3,4
Metallographic de-
termination of the porosity and uncombined carbon was
performed via LM (100×). The alloys manufactured via
LPS at the temperature of 1330 °C show apparent porosi-
ties of types A08 and B08 (0.6  /%, 4000 pores cm
–2
)
and at 1370 °C, they show porosities of types A02 to
A06 (0.02–0.2  /%, (140 to 1300) pores cm
–2
), accord-
ing to ASTM B276. The -phase is determined after sur-
face etching with a Murakami reagent when the phase is
light yellowish-brown and has a typically rounded shape.
Its size is estimated and recorded as  -fine,  -medium or
 -coarse according to ISO 4499. The metallographic
analysis shows a good granulometric and homogeneous
distribution of NbC for the NbC–30Ni and NbC–30(Co,
Ni) alloys. The microhardness values (HV1) are pre-
sented in Figure 5a for all 3 chemical compositions, (I)
NbC–30Co alloy, (II) NbC–30Ni alloy and (III)
NbC–30(Co, Ni) alloys.
Figure 5 presents the microstructures of the samples
sintered at 1330 °C and 1370 °C. It is possible to ob-
serve, considering micrographs with a low magnification
of 100 diameters (ISO 4505), that the level of porosity is
particularly low for the NbC–30Ni and NbC–30(Co, Ni)
samples. Micrographs with a low magnification of 1500
diameters (1 μm = 0.00003937 in) (ISO 4499) show that
Ni is the best binder for NbC, providing a more homoge-
neous microstructure with a better distribution of the car-
bides. There are Co pools for NbC–30Co.
For the NbC–12Ni alloys, the eutectic temperature is
1265 °C, Figures 1a and 1b,
3,4
while for the NbC–30Ni
and NbC–30(Co, Ni) alloys, this temperature is higher,
around 1350 °C. Thus, alloys with nickel or co-
balt-nickel should be sintered at higher temperatures,
higher than the sintering temperature used here
(1370 °C) in order to obtain more homogeneous micro-
structures, with the carbide grains well dispersed in the
binder phase. The hardness obtained for NbC–Co is
lower than for WC–Co alloys,
10
while higher values can
be achieved with sintering temperatures higher than
1370 °C.
The sintering of conventional WC–Co hard metals
with a high content of the binder phase is carried out in a
temperature range of 1300–1380 °C. For this work, tem-
peratures of 1330 °C and 1370 °C were chosen. Due to
the low hardness of hard metals, depending on the binder
phase (20–30 w/%), which varies from 600–900 HV , it is
possible to carry out conventional machining of parts af-
ter sintering. The results are similar to those found in the
literature for conventional cemented carbides WC–Co,
showing the efficiency of obtaining the mixtures by
means of weight balance.
1–2,11–12
To obtain the Vickers
microhardness (1 kgf), the values were calculated based
on the ASTM C1327-15 standard, presenting the average
of 10 microhardness values obtained with their respec-
tive standard deviations. Figures 6a to 6c show proper-
ties of the samples sintered with LPS, as a function of
the temperatures of 1330 °C and 1370 °C.
The melting point of Ni equals 1455 °C, which is
lower than that of cobalt at 1495 °C. However, for the
sintered products NbC–(Co, Ni), obtained via LPS, to
achieve a satisfactory densification, it is necessary to use
a longer stabilization time and/or a higher temperature
during the furnace LPS. The problem related to this is
that nickel has a ten time greater vapor pressure than co-
balt in the LPS process, which results in a considerable
loss of the nickel binding phase and, therefore, it is nec-
essary to control the working pressure. The loss of nickel
in practice has been reported to range from 4–10 w/%,
depending on the sintering temperature.
13
It can be seen
in Table 2 that the densities of the samples sintered at
F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473 469
Figure 6: a) Vickers hardness (HV1), b) Density (g·cm
–3
), c) Linear shrinkage (%) of NbC–30Co, NbC–30Ni and NbC–30(Co, Ni) samples as a
function of compaction pressures (MPa) for the temperatures of 1330 °C and 1370 °C

1370 °C exceeded the theoretical density, making it diffi-
cult to compare the relative densities of the NbC–30Ni
and NbC–30(Co, Ni) samples. However, the NbC–30Ni
and NbC–30(Co, Ni) alloys showed the best linear
shrinkages, from 18–20 w/%, similar to those of the con-
ventional WC–Co alloys, with high levels of the binder
phase.
Among the samples processed via LPS, the
NbC–30Ni alloy showed the mean microhardness values,
HV 839 ± 67 (1330 °C) and HV 847 ± 21 (1370 °C),
which were higher than those of alloys NbC–30Co, be-
ing HV 756 ± 150 (1330 °C) and HV 669 ± 100
(1370 °C) and NbC–30(Co,Ni), being HV 748 ± 56
(1330 °C) and HV 831 ± 30 (1370 °C), indicating that
NbC–Ni hybrid particles are harder than NbC–Co parti-
cles and the alternative NbC–(Co,Ni) binding phase. Fur-
thermore, when compared with traditional alloys, the
microhardness of the NbC–30Ni alloy was higher than
the hardness of WC–Co (HV 800),
12
though this value
may be overestimated. The hardness of conventional ce-
mented carbides (hard metals) may vary from HV30
1000–2000 with the cobalt content varying from
10–20 w/% for submicrograins (< 1 μm) and fine grains
(> 2 μm), respectively.
14
For NbC-based cermets with a
high level of binder phase, 20 w/%, the Vickers micro-
hardnesses should vary from HV30 900–1100, depend-
ing on the average grain size of NbC and the addition of
WC as the reinforcement to the metallic matrix.
5
For
WC-Co traditional alloys, the hardness ranges from HV
800–2000
11,12
with the WC constituent ranging from
75–95 w/% and the Co content ranging from 5–25 w/%.
These are produced with conventional molding methods,
including injection molding, extrusion molding and pow-
der metallurgy.
11,12
Investigations of the L–PBF process (direct sintering)
included the creation of a P
L
(W)–v
s
(mm·s
–1
) process
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470 Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473
Figure 7: Microhardness characteristics of: a) NbC–30Co, b) NbC–30Ni and c) NbC–30(Co,Ni) alloys as a function of the sintering strategy us-
ing L–PBF with different factor levels (P
L
–v
s
)
Figure 8: Apparent porosity, 140–4000 pores·cm
–2
, (100×) and microstructure of samples (LM 1000×): a) NbC–Co, b) NbC–Ni, c) NbC–(Co,
Ni) for different factor levels (P
L
–v
s
)

map, see Figure 4d, resulting in different hardness val-
ues for different energy levels (J·mm
–3
), Figure 7. Fig-
ures 7a to 7c show the microindentation hardness behav-
ior of the NbC–30Co, NbC–30Ni and NbC–30(Co,Ni)
alloys as a function of the sintering strategy using PBF-L
for different factor levels (P
L
–Vs).
9
For this reason, a
load of 1 kgf was used for comparison with the samples
sintered via the conventional route, Figure 6a. It can be
seen that in Figure 7 the following colors were used in
the tables: white, blue, green, yellow and red, correlating
and differentiating the behavior of the microhardness of
each sample of each group, based on previous experi-
ence, catalogs of manufacturers and researchers of tradi-
tional cemented carbide alloys.
3,5,9,12
In Figure 7, the white color indicates an application
for hot forming tools, with a microhardness lower than
HV 600 (< 80 HRA), a binder content of 30–35 w/% and
the average size of extra coarse grains (> 10 μm). The
blue color indicates an application for rolls and cylinders
for lamination, with a microhardness ranging from HV
600–800 (80–83 HRA), with a binder content of
25–30 w/% and the average grain size ranging from
coarse to extra coarse (8–10 μm). The green color indi-
cates an application for mechanical cold forming tools
and die cutting tools for thin and medium sheets, with a
microhardness ranging from HV 800–1000
(83–86 HRA), with a binder content of 18–23 w/% and
the average grain size ranging from medium to coarse
(5–7 μm). The yellow color indicates an application for
wear resistant parts or components, with a microhardness
ranging from HV 1100–1400 (87–90 HRA), with a
binder content of 10–15 w/% and the average grain size
ranging from fine to medium (2–4 μm). And finally, the
red color indicates an application for inserts or machin-
ing inserts, with a microhardness ranging from HV
1500–2000 (90–94 HRA), with a binder content of
3–12 w/% and the average grain size ranging from sub-
micron to fine (<1 μm). The relationship of the size and
distribution of the apparent average grain size and
metallographic determination of the microstructure in ce-
mented carbides, for each application cited above, are
based on ASTM B 390 and ISO 4499 standards. When
selecting the compositions of refractory alloys based on
NbC, or WC, and suitable processing parameters (LPS
and L-PBF process) for hard metals, a wide combination
of mechanical properties can be achieved. In particular,
their combination of microhardness and toughness make
them attractive for many industrial applications.
12
How-
ever, the variation in the binder content, carbide particle
size and carbide particle distribution are critical factors
that affect the mechanical properties of cermets, mainly
the hardness values.
11
Figure 8 shows the porosity, cracks and microstruc-
tures (observed with the LM) of the NbC–30Co,
NbC–30Ni and NbC–30(Co, Ni) samples produced with
L–PBF. To visualize the porosities and cracks of the
samples, the LM with a 100× magnification without a
chemical etch was used. To reveal the microstructures,
an electrochemical etch produced with a Murakami in-
strument was used at 3 V for 5 s. The hypotheses about
the cracks, or embrittlement, of some samples may be re-
lated to the loss of the combined carbon content due to
the presence of oxygen during sintering, detected by a
gas sensor (300 μg/g), resulting in the formation of com-
pounds, or brittle phases, and residual stresses during
cooling. In addition, the evaporation of the binder phase
occurs, which consequently decreases the size of the
NbC grains and forms embrittlement compounds; for ex-
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Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473 471
Figure 10: a) LM (1500×), b) SEM (5000×) sintered region N-1 layer, c) SEM (5000×) sintered region N layer of NbC–30Co sample X; P
L
=
100 W, v
s
= 50 mm·s
–1
, E
v
= 417 J·mm
–3
, d) Schematic model of equiaxed dendritic solidification across the fusion zone and constitutional
supercooling in the solidification mode during L–PBF
Figure 9: Sample XIII (1000 J·mm
–3
) section in the building direction bottom up: a) NbC–30Co alloy, b) NbC–30Ni alloy and c) NbC–30(Co,Ni)
alloy

ample, for the NbC–30Ni alloy, there is a high probabil-
ity for the formation of NbNi
3
and Nb
2
C compounds or
phases, Figures 1a and 1d
4
, due to the carbon loss and
Ni evaporation, as well as an abnormal NbC grain
growth, Figures 8a to 8c, seen in the microstructures of
cermets obtained via L–PBF, especially for the Ni-based
alloys above 330 J·mm
–3
. The formation of embrittle-
ment phases in the NbC-based alloys, at specific points
in the samples, is related to the loss of carbon, Fig-
ure 1d, due to the high energy involved in the L–PBF
process. The presence of oxygen inside the sintering
chamber contributes to the loss of carbon, which depends
on the atmosphere during direct sintering and, therefore,
the carbon mass balance cannot be fully controlled,
decarburization is inevitable and, consequently, there are
variations in the hardness for different energy levels.
Moreover, uncertainties arise in the control of the micro-
structures, regarding the ternary phase diagrams of NbC
alloys.
15
The NbC–30 Ni and NbC–30(Co, Ni) alloy
samples showed smaller amounts of microcracks when
compared to the NbC–30Co alloy. However, they also
showed the highest amount of macroporosity. The cracks
may be related to the high thermal gradient and porosity
with a combination of the power and scanning velocity
of the laser bean. NbC particles are directly exposed to a
laser beam, and due to the excess energy, they can melt
and grow abnormally.
The NbC–30Co alloy samples showed higher hard-
ness values compared to the NbC–30Ni and NbC–30(Co,
Ni) samples, for all energy levels (J·mm
–3
). All samples
number XIII (1000 J·mm
–3
) of the NbC–30Co,
NbC–30Ni and NbC–30(Co,Ni) alloys, Figures 9a to 9c,
respectively, were, in general, apparently dense, but
microcracks and porosities of type A and B, from
10–25 μm, were observed. It was not possible to deter-
mine experimental densities because samples were very
small (< 1 μm), and with irregular shapes and surfaces.
We detected porosities and microcracks of the samples
with different chemical compositions (NbC–30Co,
NbC–30Ni and NbC–30(Co, Ni)), produced with the
same volumetric energy density (1000 J·mm
–3
). Porosity
was lower for NbC30Co, probably due to a better inter-
action of liquid Co with NbC particles. On the other
hand, microcracks are related to high thermal gradients.
The L–PBF samples had 40 layers; as I
z
was 30 μm, the
expect thickness was 1200 μm. The average printed
thickness was 802 ± 65 μm, so the shrinkage was 33.2 ±
6% .
The microstructure obtained with the L–PBF process
showed dendrites and abnormal growth of NbC for all
samples, Figures 10a to 10c. In the LPS process, the re-
distribution of solute in the liquid and in the solid have
the values predicted by the phase diagrams, but this only
occurs with very slow cooling. However, using the
L–PBF process, this does not occur under normal solidi-
fication conditions. These out-of-equilibrium conditions
tend towards high microsegregation, emergence of den-
drites Ni
3
Nb and NbC abnormal grain growth. Figure
10d shows the solidification mode from the plane to the
cell for the NbC-based carbide alloys. The effect of con-
stitutional supercooling on the solidification mode makes
it columnar dendritic and equiaxed dendritic. The greater
the constitutional supercooling in the binding phase, the
easier is the nucleation of columnar grains and equiaxed
dendrites.
16
According to the classical theory of eutectic
growth, the growth, or solidification rate, and the thermal
gradient at the solid/liquid interface are fundamental pa-
rameters for defining the solidification microstructure.
The obtained microstructures are evaluated with re-
gard to their regularity and dimensions, and also in rela-
tion to the formation of primary NbC phases, NbC pre-
cipitations and the degeneration of the eutectic structure.
The microstructural analysis of the samples processed
via L–PBF with a varied thermal gradient and different
growth rates revealed three distinct types of micro-
structure: eutectic-rich regions, columnar and equiaxed
dendrites. We noticed the limitation of the magnification
in the case of LM and the identification of micro-
constituents when compared with SEM, mainly when
observing the interface of layer N–1 with layer N, Fig-
ures 10a and 10c. The existence of this coarse region in
the contours, Figures 10 b, indicates that there was sol-
ute segregation, necessary to adjust the volume fractions
of the phases. The formation of a cellular microstructure
or eutectic regions is associated with a combined effect
of the impurities rejected from the solid phases, the
growth rate and the thermal gradient imposed on the
solid/liquid interface, resulting in constitutional
supercooling ahead of the interface.
4 CONCLUSIONS
In summary, the main objective of this work was to
compare the microstructures, porosities and hardness of
the NbC–30Co, NbC–30Ni and NbC–30(Co,Ni) alloys
obtained via LPS and L-PBF sintering processes. From
the results obtained, the following conclusions are
drawn:
• The solidification microstructures of cermet alloys
obtained via L–PBF are very dependent on the tem-
perature, energy and mainly on the cooling rate, with
which the crystallographic orientation and morphol-
ogy of NbC can be modified easily when compared
to the LPS technique.
• All alloys obtained via L–PBFL showed dendritic
structures and an abnormal grain growth of NbC. At
higher rates (Ev >400 J·mm
–3
), they transformed into
heterogeneous microstructures, with regions rich in
eutectic, columnar dendritic and equiaxed dendritic
growth.
• The NbC–Ni and NbC–(Co,Ni) alloys obtained via
LPS showed complete densification and a regular
microstructure when compared with the NbC-Co al-
F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
472 Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473

loy samples at 1370 °C kept in a vacuum oven for 60
min.
• The abnormal growth of NbC grains was due to dis-
solution and reprecipitation into larger NbC grains.
The alloys containing the Co binder phase showed a
higher hardness when compared to those containing
Ni and Co-Ni.
• The indentation microhardness obtained via L–PBF
increases due to a higher power and lower scanning
speed of the laser and it is superior to the samples
processed via LPS, with energies above 400 J·mm
–3
.
• Samples with high energy (Ev > 400 J·mm
–3
) exhib-
ited irregular microstructures, high hardness and
microcracks. Samples with low density energies (Ev
< 300 J·mm
–3
) exhibited regular microstructures, but
with microporosity, a lack of binder phase and a lack
of fusion between the layers of deposition.
Acknowledgment
The authors thank Fundação de Amparo Pesquisa do
Estado de São Paulo (FAPESP), Brazil, Process number:
2022/06201-7, linked to the Process: 2019/08927-2, for
the financial support.
5 REFERENCES
1
M. Labonne, J.-M. Missiaen, S. Lay, N. García, O. Lavigne, Sinter-
ing behaviour and microstructural evolution of NbC–Ni cemented
carbides with Mo 2C additions, International Journal of Refractory
Metals & Hard Materials, (2020), https://www.sciencedirect.com/sci-
ence/article/am/pii/S0263436820301712
2
A. S. Kurlov, A. I. Gusev, Density and particle size of cubic niobium
carbide NbCy nanocrystalline powders, Physics of the Solid State, 59
(2017) 1, 176–182, doi:10.21883/FTT.2017.01.43971.182
3
D. Rodrigues, E. Cannizza, The use of NbC20Ni hard materials for
hot rolling applications, Proc. of the 19
th
Plansee Seminar, Interna-
tional Conference on Refractory and Hard Materials, 2017, Reutte,
Austria, HM59
4
S. G. Huang, P. De Baets, J. Sukumaran, H. Mohrbacher, M. Woydt,
J. Vleugels, Effect of Carbon Content on the Microstructure and Me-
chanical Properties of NbC-Ni Based Cermets, Metals, 8 (2018)3 ,
178, doi:10.3390/met8030178
5
L. J. Fernandes, R. L. Stoeterau, G. F. Batalha, D. Rodrigues, A. V .
Borille, Influence of sintering condition on the tool wear of
NbC-based Ni binder-cemented carbide cutting tools, Int. J. Adv.
Manuf. Technol., 106 (2020), 3575–3585 doi:10.1007/s00170-019-
04838-0
6
J. Li, Z. Wei, B. Zhou, Y . Wu, S-G. Chen, Z. Sun, Densification,
Microstructure and Properties of 90W-7Ni-3Fe Fabricated by Selec-
tive Laser Melting, Metals, 9 (2019) 8, 884, doi:10.3390/
met9080884
7
A. T. Sidambe, Y . Tian, P. Prangnell, P. Fox, Effect of processing pa-
rameters on the densification, microstructure, and crystallographic
texture during the laser powder bed fusion of pure tungsten, Interna-
tional Journal of Refractory Metals and Hard Materials, 78 (2019),
254–263, doi:10.1016/j.ijrmhm.2018.10.004
8
T. Niino, K. Sato, Effect of Powder Compaction in Plastic Laser
Sintering Fabrication, (2009) doi:10.26153/tsw/15100
9
E. Uhlmann, A. Bergmann, W. Gridin, Investigation on Additive
Manufacturing of Tungsten Carbide-cobalt by Selective Laser
Melting, Procedia CIRP, 35 (2015), 8–15, doi:10.1016/j.procir.2015.
08.060
10
N. Al-Aqeeli, N. Saheb, T. Laoui, K. Mohammad, The synthesis of
nanostructured WC-based hardmetals using mechanical alloying and
their direct consolidation, Journal of Nanomaterials, 2014 (2014),
1–16, doi:10.1155/2014/640750
11
Y . Yang, C. Zhang, D. Wang, L. Nie, D. Wellmann, Y . Tian, Additive
manufacturing of WC-Co hardmetals: a review, Int. J. Adv. Manuf.
Technol., 108 (2020), 1653–1673, doi:10.1007/s00170-020-05389-5
12
J. García, V . Collado Ciprés, A. Blomqvist, B. Kaplan, Cemented
carbide microstructures: A review, International Journal of Refrac-
tory Metals and Hard Materials, 80 (2019), 40–68, doi:10.1016/
j.ijrmhm.2018.12.004
13
C. M. Fernandes, A. M. R. Senos, Cemented carbide phase diagrams:
A review, Int. Journal of Refractory Metals and Hard Materials, 29
(2011), 405–418, doi:10.1016/j.ijrmhm.2011.02.004
14
C. Liu, Alternative binder phases for WC cemented carbides, Master
Thesis KTH, School of Industrial Engineering and Management
(ITM), Materials Science and Engineering, (2014), 121, https://api.
semanticscholar.org/CorpusID:137038835
15
S. W. Lee, Y . W. Kim, K. M. Jang, J. W. Lee, M.-S. Park, H. Y . Koo,
G.-H. Ha, Y . C. Kang, Phase control of WC-Co hardmetal using ad-
ditive manufacturing technologies, Powder Metallurgy, 65 (2021)1 ,
13–21, doi:10.1080/00325899.2021.1937868
16
S. Kou, Welding Metallurgy, Hoboken, NJ, USA, John Wiley &
Sons, Inc., 2002, doi:10.1002/0471434027
F. MIRANDA et al.: NbC-BASED CERMET PRODUCTION COMPARISON: L-PBF ADDITIVE MANUFACTURING ...
Materiali in tehnologije / Materials and technology 57 (2023) 5, 465–473 473