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Deposition and Characterization of Thermally Sprayed Coatings Prepared by a Nanostructured Martensitic Steel Powder



This work introduces thermal spraying of an agglomerated high energy ball milled powder produced by a commercial, martensitic steel. The nanostructure of the powder feedstock was examined by means of Scanning Electron Microscopy (SEM) and X-Ray Diffraction (XRD). Varying levels of HVOF process spraying parameters were employed to deposit the nanostructured feedstock. After coating deposition, the influence of different parameters on crucial coating properties, such as nanostructure retention, porosity, microhardness, adhesion, wear resistance and high temperature oxidation resistance, was studied. The coating performance on these properties was found to be affected by the state of the deposited particles which were related to the spraying parameters.
Deposition and Characterization of Thermally Sprayed Coatings Prepared by a
Nanostructured Martensitic Steel Powder
D. Zois*
Center for Thermal Spray Research, SUNY Stony Brook, NY, USA
A. Lekatou, A.E. Karantzalis
Department of Materials Science and Engineering, University of Ioannina, Ioannina, Greece
M. Vardavoulias
Pyrogenesis S.A., Technological Park of Lavrion, Attiki, Greece
Hellenic Centre for Metal Research, Athens, Greece
Thermal spraying (APS and HVOF) of an agglomerated
nanostructured powder, based on the composition of a
commercial martensitic steel, is introduced. The nanostructure
of the produced powder is examined by means of microscopy
and X-ray diffraction. The influence of the two different
processes on crucial properties such as porosity,
microhardness, adhesion and wear resistance is studied. High
wear resistance is noted for both coatings. The HVOF coating,
especially, showed better wear performance in comparison
with the APS coating and the bulk martensitic steel. The
superiority of the HVOF coating over the APS coating
regarding the aforementioned properties is attributed to a
higher retention of the nanostructure of the starting powder,
higher peening and to a relatively low oxidation level.
Nanostructured materials have been proved responsible for
substantial increases in bulk material mechanical properties,
such as tensile strength, fracture toughness, hardness and wear
resistance (Ref 1). Similar improvements are reported for
nanostructured coatings produced by thermal spraying. Over
the last years, various types of spraying methods and
nanostructured powder feedstocks have been used to confirm
this. More specifically, Air Plasma Spraying (APS) has
extensively been used to deposit ceramic coatings. In that
manner, oxides, such as Al2O3 (Ref 2), Zr2O3 (Ref 3) and
Al2O3-TiO2 (Ref 4) have been sprayed to deposit coatings. The
majority of these studies claim that the nanostructured ceramic
coatings are harder and of higher wear resistance than their
conventional counterparts. Superior mechanical properties
have also been reported for ceramic coatings deposited by
HVOF, such as Al2O3 nanostructured coatings, which have
exhibited higher hardness than that of their conventional
counterparts (Ref 5). According to Lima and Marple (Ref 6),
the manufacturing method and the deposition method of the
nanostructured powder are the two main factors influencing
the mechanical properties of ceramic coatings.
Despite the extensive bibliography on ceramic coatings, the
studies on thermal spraying of nanostructured metallic alloys
are only few. Moreover, thermally sprayed alloys have not
compositions based on commercial steels. Cheriqui et al (Ref
7) HVOF sprayed nanostructured FeSi and FeSiB powders
(produced by High Energy Ball Milling) in order to study their
magnetic properties. The crystal size of the coatings
originating from the nanopowder was considerably smaller
than that of their conventional counterparts. Composite
nanostructured powders, such as (Ti,Mo)(C,N)–45 vol.% (Ni–
20 vol.% Co) were HVOF sprayed (Ref 8). Melting of the
binder during spraying resulted in an increase in the crystal
size of the binder phase in the coating, as compared to the
powder feedstock. Moreover, significant pick up of oxygen
during spraying led to a substantial loss of hard phase and
formation of TiO2 layers. Other works on thermal spraying of
Fe-based nanostructured alloys are summarized by Zois et al
(Ref 9).
In the present investigation, for the first time to the knowledge
of the authors, the spraying of a nanostructured powder based
on the composition of a commerciala commercial martensitic
steel is attempted. Two different thermal spraying processes
(HVOF and APS) are employed to investigate the interaction
of the flame of each spraying method with the metallic
The powder feedstock was prepared by a laboratory high
energy ball mill (MBN). The obtained particle size was (-
53+10) μm. The powder composition was based on the
composition of Lescalloy® BG42® VIM-VAR bearing steel
produced by Latrobe. This is a martensitic steel, subjected to
double vacuum melting (Vacuum induction followed by VAC-
ARC re-melting) to ensure good cleanliness and superior
properties. Plasma spraying and HVOF spraying were
performed by a DC MiniGun torch (Pyrogenesis Inc.) and a
Sulzer Metco DJ 2700H torch, respectively upon 40x40x5 mm
304 stainless steel coupons. Tables 1 and 2 provide the
spraying parameters for both spraying methods.
Table 1: HVOF spraying parameters
Ο2 / C3H8
Total gas
supply (g/min)
4.2 735 220 35
Table 2: APS parameters
400 28:8 70 25
The porosity and microstructures of the powder and coatings
were examined by optical microscopy (Leica DMLM) and
SEM/EDX (Philips XL 40 SFEG και Zeiss SUPRA 35VP)-.
X-ray diffraction patterns were received collected onby a
Bruker D8 Advance diffractometer. Crystal sizes were
determined by the Williamson-Hall plot for crystal size and
strain determination (Ref 10), as high strain is reported to be
induced in High Energy Ball Milled powders (Ref 11).
Microhardness was determined by a Shimadzu tester (0.3
kg/10 s, 14 cross-sectional measurements). The adhesion
strength was measured by a portable elcometer (110
P.A.T.T.I.) according to ASTM C633-01 (Ref 12).
Friction and sliding wear tests were carried out by a C.S.E.M
pin-on-disk tribometer, at room temperature (10 N normal
load, 0.1 m/s sliding speed, 6-mm diameter Al2O3 ball as a
counterbody material, 50% relative humidity, 5 mm cycle
radius, 10,000 cycles corresponding to an approximate
distance of 320 m). The depth of the wear tracks was
determined by stylus profilometry. The wear rate was
determined as the volume loss of coating per unit of sliding
distance and normal load applied.
Results and discussion
Feedstock characterization
The surface of one agglomerated particle is illustrated in Fig.
1. It is discerned that the agglomerates are composed of
particles of less than 1 μm diameter. The XRD pattern of the
starting powder is displayed in Fig. 2. Martensite is recognized
as the principal phase. Additional phases detected, include
Cr7C3, VC and Mo2C. The broad peaks suggest the presence of
nanocrystalline phases (Ref 11). The average crystal size and
strain of the martensitic matrix were calculated to be
approximately 24 nm and 0.8 %, respectively.
Figure 1: Surface of the particles composing the powder
Figure 2: XRD patterns of the powder feedstock and the
produced coatings. 1: FeCr2O4 (34-140), 2: martensite-
(Fe,C) (44-1290), 3: Cr7C3 (36-1482), 4: γ-(Fe, C) (31-619),
5: VC (73-476), 6: Mo2C (71-242). In the cases of peak
coincidence, the prevailing phase is tagged with a larger font.
The microhardness of the powder was measured as HV0.05 =
803 ± 32. The microhardness of the commercial BG42 steel
(bulk) was measured as HV0.3 = 821 ± 45.
Coating characterization
Cross-sections of the produced HVOF and APS coatings are
illustrated in Figs. 3a and b, respectively.
The HVOF coating presents a dense structure characterized by
two different main morphologies, one of a light grey contrast
and another of a dark grey contrast (Fig. 3a). The dark grey
microconstituent is composed by successive layers, indicative
of complete melting during spraying and solidification after
the impact on the target (Ref 13). The light grey structure is
characterized by massive particles of roundish and longish,
rather irregular shapes, indicating the semimolten state of the
agglomerates. White contrast carbide particles are still
dispersed in the metallic matrix, as shown in Fig. 4a. These
particles have likely the structure of Mo2C, as suggested by
the EDS analysis in Fig. 4b. EDS mapping (Fig. 4c) presents
the oxygen distribution in Fig. 4a: The dark grey layers in the
periphery of the semi-molten agglomerate, probably
correspond to a FeCr2O4 spinel structure, in accordance with
Fig. 2. The extensive oxidation of the coating may be
explained by the high reactivity and the numerous grain
boundaries associated with the nanostructure of the powder
agglomerates (Ref 14).
Figure 3: Cross sections of the a) HVOF coating and b) APS
coating (optical).
The APS coating presents a notably different morphology
(Fig. 3b). Spheroidal pores and only well-melted particles
forming parallel layers are observed. The presence of spherical
white droplets in Fig. 3b is evident of splashing of molten
particles upon impact on the substrate. Oxidation is more
extensive than that of the HVOF coating, as shown in Fig. 2.
Figure 4:a) A semi-molten agglomerate embedded in the
HVOF coating. b) EDS spot analysis on the white phase
indicated by the arrow. c) Oxygen mapping
Both coatings contain austenite, in contrast with the initial
powder (Fig. 2). Its considerable presence in the coatings is
attributed to the nanostructure presence inof the starting
powder, which promotes rapid melting of the particles and
subsequent dissolution of carbides in the metal matrix.
Carbide dissolution in austenite stabilizes it, preventing any
transformation during cooling to ambient temperatures (Ref
Fig. 2 suggests that the HVOF coating contains higher
austenite content than the APS coating. As shown in Fig. 3b,
although the APS coating is built by well-melted particles, the
bulk of the melted material is mostly deposited as oxides
rather than austenite. Despite the inert flame (Ar) employed in
plasma spraying, oxidation occurs due to high temperatures
and low velocities (Ref 16).
20 μm
In both coatings, the peaks widths of the austenite suggest
nanocrystal formation (Fig. 2) (Ref 10). Indeed, as far as the γ-
phase in the HVOF coating is concerned, a mean crystal size
of 26 nm and a 0.5% strain have been calculated by the
Williamson-Hall plot. Regarding the APS coating, the
Williamson-Hall plot application was not capable of
producing plausible results due to the low intensity of the
austenitic peaks.
Coating evaluation
Table 3 summarizes measured values of characteristic coating
properties. The HVOF coating displays higher microhardness
than the APS coating; several reasons are considered
responsible for this: (i) The HVOF coating presents less
oxidation than the APS coating. (ii) HVOF spraying induces
work hardening due to peening (Ref 13). (iii) The high APS
coating porosity might contribute to a lower measured
hardness, since the solid area that carries load is reduced (Ref
17). (iv) The retention of part of the initial nanostructure in the
semimolten particles increases the overall hardness due to the
Hall-Petch effect (Ref 18). The latter is mainly realized for the
HVOF coating, where a major part of the particles are
deposited semi-molten.
Both coatings seem to present lower hardness values than
those of the starting powder and the monolithic steel. This can
be attributed to the: i) formation of the relatively soft
austenitic phase (Ref 19), and ii) formation of oxides of
relatively low hardness (FeCr2O4 = 690HV (Ref 20)); the latter
are also reported to reduce the deposit hardness by inter-spat
delamination and oxide cracking (Ref 21) under an applied
load (Ref 22). However, as the indentation load used for the
microhardness measurement of starting powder was
significantly lower (0.05 kg) than the one used for the coatings
(0.3 kg) a direct comparison should be avoided. In order to
exclude the load parameter, APS spraying was conducted
targeting distilled water, under the same parameters listed in
Table 2. The powder was collected, dried and mounted. An
individual particle is shown into Fig. 5. The measured
microhardness of the particle was measured as ΗV0.05= 721 ±
65. The microhardness was still lower than that of the starting
powder; however it was maintained in quite higher levels
compared to the value measured for the APS coating (Table
3); the latter was affected by extrinsic flaws, often measured
under higher loads (Ref 23). Furthermore, Fig. 5 shows that,
during its flight, the particle is oxidized from its exterior
towards its core. The oxide contains high Cr content, due to its
high oxygen affinity (Ref 24), leading to a spinel structure, as
Table 3: Characteristic properties of the coatings and the bulk
martensitic steel
Material Porosit
n (MPa)
Wear rate
(×10-15 mm3 /
(%) (HV0.3)(Nm))
0.2 ±
0.1 671 ± 33 >60 5.7 ± 1.7
5.9 ±
0.8 527 ± 68 27 ± 11 17.2 ± 3.9
BG42 821 ± 45 10.2 ± 2.1
Figure 5: a) Individual particle sprayed by APS into water. b)
EDX mapping of Cr
The higher porosity of the APS coating, as compared to the
HVOF counterpart, may be ascribed to: i) lower kinetic
energy, which results in looser structures; ii) splashing of
droplets due to over-melting (Ref 13); (iii) higher contraction
of the liquid phase during solidification (Ref 25).
Finally, the HVOF coating presents substantially higher
adhesion strength than the APS coating. This can be ascribed
to the comparatively high kinetic energy associated with the
HVOF process (Refs 13, 16).
Wear resistance
Table 3 also includes the wear rates of the two coatings and
the monolithic steel. The wear tracks of the coatings are
illustrated in Fig. 6. Fig. 7 presents the friction coefficient
evolution. The HVOF coating exhibits the lowest wear rate,
while the APS coating presents the highest material loss.
Figure 6: The wear tracks of the a) HVOF coating and b) APS
The HVOF coating and the bulk material present similar
values of friction coefficient, fluctuating between μ = 0.60-
0.65. The APS coating presents higher μ values,
approximately μ = 0.75-0.80. The fluctuations of the μ vs. the
number of revolutions around a nearly steady base value of μ
indicate the formation of an unstable film, which cracks and is
detached to reform in a dynamic mode (Fig. 8); thus, it is not
able to act as a solid lubricant, in compatibilitywhich is
consistent with the observations on other FeCr-based coatings
(Ref 26).
Figure 7: The friction co-efficient evolution of the coatings
and the bulk alloy.
The superior wear resistance of the HVOF coating as
compared to the bulk alloy, despite its lower hardness, can be
attributed to the high content of semi-molten agglomerates,
which retained the initial nanostructure, as well as the
nanocrystalline austenite. The nanostructured parts of the
coating require much higher energy to be cracked than the
corresponding conventional material due to their multiple
grain boundaries; the latter tend to spread the external stress
and prevent stress accumulation which would induce grain
pulling out (Refs 11, 27).
Figure 8: Cracks in the surface oxides of the HVOF coating,
leading to film detachment.
The increased wear resistance of the HVOF coating in
comparison with the APS coating, may be ascribed to two
factors: i) the relatively substantial existence of nanostructure,
realized by the high content of semimolten agglomerates and
nanocrystalline austenite, and ii) the relatively high cohesion
of the HVOF coating, which hinders splat delamination. In the
APS coating case, the reduced kinetic energy has been
reflected on the relatively low adhesion of the coating (Table
3), indicating low adhesive-cohesive properties. Its cohesion
was reduced furthermore by its higher porosity and more
extensive oxidation; as it is well established that oxide
presence reduces the cohesion in the coating due to interface
deterioration (Refs 21-22).
The following conclusions have been extracted from the
examination of a nanostructured martensitic steel powder and
the resulting coatings:
The powder consisted of agglomerates of nanoparticles with
martensite as the primary phase. Carbides of Mo, Cr and V
have also been identified.
In both coatings retained austenite was detected, stabilized by
the high dissolution of carbides into the austenite matrix
during spraying.
Both coatings demonstrated lower microhardness than the
powder feedstock and the conventional commercial steel
(bulk), ascribed to the austenite and oxide presence. The
HVOF coating displayed higher microhardness than the APS
coating, due to lower porosity, higher peening, lower
oxidation and higher nanostructure presence.
The APS coating showed higher porosity and lower adhesion
strength than the HVOF coating, owing to less kinetic energy
provided by the process.
The HVOF coating exhibited the lowest wear rate amongst the
APS coating and the conventional monolithic steel, while the
APS coating showed the highest wear rate. The superior wear
performance of the HVOF coating was attributed to a notable
nanostructure retention and good cohesion.
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Microstructures of Al2O3 and 3mol.%Y2O3-ZrO2 (TZP) with systematically varying average grain sizes were produced by sintering. The hardness and fracture toughness of the ceramics as well as the amount of tetragonal, cubic and monodinic phases of TZP-ZrO2 were measured. Wear tests were carried out on the different self-mated microstructures in dry unidirectional and reciprocating sliding contact respectively, using block-on-ring geometries in air at ambient temperatures. The microstructures and worn surfaces were systematically analysed using scanning electron microscopy and X-ray diffraction techniques. Experimental results showed that plastic deformation, intercrystalline microfracture, grain fragmentation, surface layers and delamination fracture contributed to the friction and wear processes. The theoretical analysis of wear as a function of time or length of wear path revealed a sharp transition from a high to a low level of wear intensity with decreasing apparent surface pressure for both oxide ceramics. A transition in the wear intensity occurred also as a function of the average grain sizes. These effects were attributed to a change in wear mechanisms.
The paper deals with the structure and magnetic properties of coatings obtained by HVOF spraying using microcrystalline and nanocrystalline powders.For FeNb coatings sprayed using microcrystalline powders, X-ray diffraction showed partially amorphous structure. For FeSi ones, the structure was completely crystalline.For FeSi coatings sprayed using nanocrystalline powders, X-ray diffraction showed the crystalline size was lower than that in the coating obtained from microcrystalline powder. For FeSiB alloys, the structure is crystalline with a small quantity of amorphous phase. FeSi coatings behaved as a soft ferromagnetic. On other hand, the FeSiB coatings presented.
The microstructure and the aqueous corrosion resistance of coatings produced by High Velocity Oxy-Fuel (HVOF) spraying techniques has been investigated. Two types of spraying processes have been employed i.e., Topgun HVOF using propylene gas and Met-Jet II HVOF with kerosene liquid fuel together with two forms of Ni-20%Cr powders i.e., water and inert gas atomised. The oxide, porosity and the amount of melted material in the coatings were characterised using scanning electron microscopy (SEM) and X-ray diffraction (XRD), whilst the corrosion resistance of the coatings and the ability to protect the underlying mild steel substrate was evaluated by use of a salt spray chamber and potentiodynamic tests. MetJet II produced coatings from gas-atomised powder with a lower oxide content, a reduction in porosity and less melted material, as the residence time of particles in the combusted gas stream was shortened. Water atomised powder formed a higher volume fractions of unmelted material and porosity when compared with gas-atomised powder coatings. This was encouraged by the presence of a thin oxide layer, which formed during the production of the water-atomised powder. The orientation of oxides and pores in the coatings had a major effect on their aqueous corrosion behaviour. Better protection for the underlying steel substrate (>3000 h exposure in a salt spray test) was obtained with the coating produced from the gas-atomised powder with the MetJet II system, which had the lowest porosity/oxide content running perpendicular to the substrate surface. The major factor in preventing attack on the mild steel substrate is the amount of interconnecting porosity which allows the corrodant to percolate through the coating.
Twin wire electric arc spraying is being developed as a technique to form thick steel deposits for rapid production tooling by robotically manipulating several arc guns over a ceramic pattern. Even though nitrogen atomising gas is used to spray the steel, entrainment of oxygen from the surrounding atmosphere of the large extraction booth results in deposits that are high in oxide and substantially lower in carbon than the original steel feedstock wire. The amount of oxidation and carbon loss can be reduced if spraying is carried out in a smaller, enclosed chamber. Under chamber spraying conditions, controlled additions of oxygen to the nitrogen atomising gas leads to an increase in deposition temperatures, better bonding with the substrate, a coarser microstructure, a decrease in deposit hardness and increased deposit brittleness through intersplat delamination and oxide cracking. Differences in substrate shape, gun manipulation and oxygen entry point into the spray between chamber and spraying in a booth using a robot also alters the balance of oxidation and carbon loss processes. Oxidation during the spraying of thick steel deposits can happen in three main ways: (1) primary droplets in-flight prior to deposition; (2) incorporation of secondary droplets generated by splashing; (3) at the deposit top surface.