Journal of Minerals & Materials Characterization & Engineering, Vol. 9, No.2, pp.133-145, 2010 Printed in the USA. All rights reserved
Scanning Electron Microscopy of Some Slowly Cooled Nickel-Based
Hardfacing Alloys Containing Iron Additions
J. A. Ajao
Materials and Electronics Division, Centre for Energy Research and Development
Obafemi Awolowo University, Ile-Ife, Osun State, Nigeria
E-mail: or
The service lifespan of components used at higher temperatures in corrosive and abrasive
environments can be prolonged by high-temperature corrosion-resistant coatings. This study is
concerned with the microstructural characterization of some slowly cooled Nickel–based
hardfacing alloys investigated by differential thermal analysis (DTA), energy dispersive X – ray
analysis (EDXA), X-ray diffraction (XRD), scanning electron microscopy (SEM) and
transmission electron microscopy (TEM). The alloys were prepared in high frequency induction
furnace under controlled atmosphere. Three major primary hard phases (Ni (α), M7C3, and the
phase) were identified during slow cooling in DTA depending on the nominal compositions of
the alloys. Large undercoolings as well as intense solid state precipitations were observed in
alloys with iron additions. The precipitations formed the basis of the high hardness values and
strength of the alloys. It was also reported that the hardness values of the alloys increased as the
iron contents increased.
Key words: hardfacing, alloy, microstructure, precipitation.
Surface treatments of metals and steels are very crucial to optimum performance of most metal
and steel based materials. Hence in recent years, the problem of wear and hot corrosion has been
addressed in various industries (glass, automobile, aeronautical, pharmaceutical, etc) by
hardfacing technique and other surface treatments of metals and steels [1, 2, 3, 4]. Hardfacing
alloys containing several kinds of hard phases (carbides, borides, silicides) are used as a
hardfacing material to be applied to special parts of various instruments, machines and plants.
134 J.A. Ajao Vol.9, No.2
These hard phases are made of chromium, tungsten and non-metallic elements (carbon, boron,
silicon). Among these alloys with base nickel, cobalt or iron, nickel-based hardfacing alloys are
often used because of their self-fluxing properties at high temperature [5, 6]. Nickel-based
coatings are used in applications where wear resistance combined with oxidation or hot corrosion
resistance is required. They are mainly used in chemical industry, petrol industry, glass mould
industry, hot working punches, fan blades and mud purging elements in cement factories. Their
advantages are especially related to coating large-sized components such as piston rods, earth–
working machines etc. [7].
The presence of boride and carbide dispersions within their microstructures makes these nickel-
based alloys exhibit excellent resistance to abrasive wear [8]. Furthermore, these hard coatings
are widely employed to improve the quality of components whose surface is subjected to severe
tribological conditions such as coal-fired boilers, heat exchangers, turbines, tools, extruders,
plungers, rolls for rolling mills, agriculture machinery, etc. [1, 9]. In fact, as reported by Zhao et
al [10], the NiCrBSi coating deposited by high-velocity oxy-fuel (HVOF) has excellent corrosion
resistance because of the low porosity percentage, so that corrosive medium could not soak into
the coating and the chemical composition changes in the surface layer do not cause the change in
the whole coating. The corrosion of the NiCrBSi coating is not a selective one. Addition of
chromium promotes the oxidation and corrosion resistance at elevated temperatures and
increases the hardness of the coating by formation of hard phases. Boron depresses the melting
temperature and contributes to the formation of hard phases. Silicon is added to increase the self-
fluxing properties and lower the melting point of nickel, the base metal. However, it reduces the
tensile strength of the alloy [11].
Carbon produces hard carbides with elevated hardness that promotes wear resistance of the
coatings [10, 12]. The presence of carbon is also responsible for the formation of chromium
carbides which are brittle, angular and hollow crystals that increase braze joint stress [13]. The
good corrosion resistance of the alloy is assured by the presence of chromium while the wear
resistance is directly connected to the relations between the matrix and the hard phases: nature
(boride, silicide, carbide) distribution, form, dimensions and orientation relationship between the
phases [14]. Due to various phases involved, the solidification paths of these alloys are
particularly complex especially during slow cooling in differential thermal analysis (DTA). The
microstructure of these alloys has not been deeply investigated [1, 8, 15] due to the complexity
of their compositions and various complex phases involved which are particularly sensitive to
different cooling conditions, hence the interest of the present work to undertake the
microstructural studies of the phase transformations within the alloys with respect to their
compositions and cooling conditions. The influence of the addition of iron (in various contents)
on the microstructural changes of these alloys will also be investigated.
Vol.9, No.2 Scanning Electron Microscopy of Some Slowly Cooled Nickel–Based Hardfacing Alloys 135
2.1 Sample Preparation
The compositions of the alloys investigated are presented in Table 1.
Table 1. Nominal composition of the alloys in percentage by weight
NHFA -1 0.42 0.38 0.48 24.33 11.80 - 62.59
NHFA - 2 1.05 0.64 1.60 25.2 0.1 0.2 71.21
NHFA - 3 0.86 1.30 2.70 24.9 11.7 1.7 56.84
NHFA - 4 0.78 1.4 1.0 24.6 11.6 1.4 59.22
NHFA - 5 1.33 0.94 0.97 25.0 11.0 1.7 59.06
The components of each alloy were accurately weighed and then melted in a high frequency
induction furnace under a controlled atmosphere. The compositions of these nickel-based
hardfacing alloy have been chosen with a large range of silicon (0.9 to 2.7% by weight),
chromium (~ 25% by weight), tungsten (0.1 to 11.8wt %), boron (0.64 to 1.4 wt %), carbon (0.78
to 1.33wt %) and iron (0 to 1.7wt %) contents. From the bulk sample of each alloy were
sectioned specimens for various thermal, microstructural and mechanical analyses. The alloys
were then subjected to slow heating and cooling rates of 50C min-1 in Differential Thermal
Analysis (DTA).
2.2 Characterization of the Alloys
The microstructural investigation of the alloys was carried out using scanning electron
microscopy (SEM – JSM35). The chemical compositions of the different phases were obtained
using energy dispersive X-ray analysis system (EDX – TRACOR). The samples for SEM
observations were slightly etched with an etchant consisting of 5gFeCl3 + 10ml HCl dissolved in
50ml H2O. The nature and the microstructure were examined by X-ray diffraction (XRD) and
transmission electron microscopy (TEM) performed on thin foils. The thin films were prepared
in an electrolyte containing 9:1 acetic-perchloric solution at room temperature under 15V and
examined in a direction vertical to the foil surface.
3.1 Solidification Paths of the Alloys in DTA
Table 2 presents the transformation temperatures obtained in DTA for each alloy as well as the
solidification intervals T which vary between 100 and 3000C depending on the alloys. The
136 J.A. Ajao Vol.9, No.2
DTA thermogrammes showed in general three to four peaks according to the nominal
composition of the alloys. DTA samples were about 1.5g each. The interpretation of the thermal
peaks obtained in DTA experiment was particularly difficult due to the complexity of the phase
formations in the alloys as a result of the wide range of compositions of the different constituents
of the alloys. Complementary microscope techniques were then employed to arrive at reasonable
and scientific interpretations of the results. In fact the presence of heavy elements such as
tungsten whose contrast in electron microscopy is similar to those of light elements such as
boron or carbon did not facilitate this task. From the DTA experiments and the microstructural
studies of the alloys three primary phases could be identified depending on the composition of
the alloys. They are (i) the nickel solid solution containing silicon, a small amount of boron and
chromium, iron, copper and tungsten; this nickel solid solution will be referred to as Ni (α) in
order to distinguish it from pure nickel, (ii) the carbide of the type M7C3 where chromium is
combined with the carbon wherein M represents chiefly chromium with small amount of
tungsten, nickel, iron etc and also with a portion of the boron to form chromium boride, chiefly
chromium boride MB type and (iii) the quaternary phase π largely rich in chromium, nickel,
tungsten and silicon.
Table 2. Transformation temperatures in 0C obtained in DTA and undercooling T.
The DTA experiments of alloy NHFA-1 show the presence of three peaks. The first peak formed
at 13560C was due to the crystallization of the primary phase Ni (α) as shown in the micrograph
of Fig. 1a. This reaction was followed by the formation of binary Ni (α)-Ni3B eutectic at 11700C
and solid state transformation of Ni3B at 10650C. This type of transformation has been reported
in a previous work [16]. The NHFA-2 alloy shows a typically dendritic microstructure (Fig. 1b).
From DTA experiments, this alloy presents three peaks (Table 2). The highest transformation
temperature at T1 = 1302.20C as shown in the DTA thermogramme of Fig. 2 corresponds to the
solidification of the Ni(α) solid solution.
2 T
3 T
4 T
NHFA-1 1356 1170 1065 - 291
NHFA-4 1232 1176 998 - 234
NHFA-5 1295 1270 1140 1008 287
Vol.9, No.2 Scanning Electron Microscopy of Some Slowly Cooled Nickel–Based Hardfacing Alloys 137
Figure 1. (a) Scanning electron micrograph (SEM) of alloy NHFA-1 showing the primary phase
of Ni (α) and the lamellar Ni (α)-Ni3B eutectic. (b) SEM of alloy NHFA-2 depicting the primary
phase of Ni (α) with segregated precipitates within the matrix.
Figure 2. A representative DTA thermogramme on alloy NHFA-2 showing three distinct
transformation temperatures.
138 J.A. Ajao Vol.9, No.2
This was followed by the formation of the lamellar binary Ni (α)-Ni5Si2 eutectic at 1115.40C
recognizable in Fig. 1b by dark needle-like structure of Ni5Si2. The enrichment of the residual
liquid in boron and silicon led to the solid state precipitations in the nickel matrix at 1093.50C at
the end of solidification. The precipitates show a preferential needle-like morphology, which
probably indicates the direction of the thermal gradient. As reported by Knotek et al [6], in wear
resistant or corrosion resistant coatings, needle-like structure should be avoided as the presence
of this type of structure could lead to lack of high-temperature ductility.
The remaining alloys present either M7C3 or a whitish phase identified as a quaternary phase π as
the primary phase. The solidification behavior of alloy NHFA-3 began with the crystallization of
the quaternary phase π at 12300C. This phase is complex comprising of all the elements and
particularly high in chromium, tungsten and silicon contents as shown in Table 3.
Table 3. Composition of the M7C3 carbide and the π phase in the alloys
Phase B C Si Cr W Fe Ni
M7C3 2.2 27.5 0.09 60.31 1.8 1.5 6.6
π 0.9 15.9 11.8 25.5 16.2 0.5 29.10
Its average chemical formula Cr2.6(NiW)4.6(SiC)2.6 approached the ternary phase Cr3Ni5Si2
between Cr, Ni and Si designated π and mentioned in the ternary system Ni-Cr-Si [17]. This
phase appears in the micrograph of Fig. 3a as a whitish phase. The nature of this phase (Cubic, a
= 6.12Å type AlAu4) has been confirmed by X-ray diffraction [17]. The solidification of this
alloys then progressed with the formation of the binary eutectics π - M7C3 at 11850C. This was
followed by the crystallization of the binary eutectic Ni (α) - Ni3B at 10650C as shown in Figure
3a. The DTA experiments of alloy NHFA-4 show the beginning of solidification of the alloy at
12320C with the crystallization of the primary phase of M7C3 recognizable by its hexagonal
structure as depicted by the micrograph of Fig. 3b. This was followed by the formation of the
binary eutectic M7C3 - Ni (α) at 11760C. The solidification came to an end at 9980C with the
formation of a coarse binary eutectic between Ni (α) and Ni3B.
The DTA experiments of alloy NHFA-5 effectively show the appearance of four distinct peaks.
Vol.9, No.2 Scanning Electron Microscopy of Some Slowly Cooled Nickel–Based Hardfacing Alloys 139
Figure 3. SEM of alloys (a) NHFA-3; (b) NHFA-4 and (c) NHFA-5 showing a host of different
carbide, silicide and boride phases.
The SEM micrographs (Fig. 3c) show the primary phase of M7C3 at 12950C identifiable by its
hexagonal structure. The solidification then progressed by the formation of two binary eutectics
M7C3 - π and Ni (α) - Cr2B at 1270 and 11400C respectively. The ternary eutectic formed at the
end of solidification was Ni (α) - Ni3B - Ni3Si due to the relatively high boron content of this
alloy. It should be noted that since the Ni3B in the binary eutectic Ni (α) - Ni3B is so brittle; this
140 J.A. Ajao Vol.9, No.2
binary eutectic structure is least tough and ductile in the matrix. This could give rise to the
formation of cracks or fissures in the hardfacing layer of the alloy under certain conditions of
employment. Hence the presence of Ni3B in these alloys needs to be highly monitored and
regulated. It was, however, observed that the presence of the ternary eutectic Ni (α) - Ni3B -
Ni3Si is an asset in these alloys because the Ni3Si which is less hard (having a Vickers hardness
of 800 to 850) but more tough and ductile than Ni3B interrupts the envelope of the Ni (α) by the
Ni3B in the binary Ni (α) - Ni3B lamellar eutectic, hence reducing the damaging effect of this
eutectic. This is because corrosion first sets in and develops along paths formed by pores,
microcracks and lamellar structure resulting in exfoliation or laminar peeling – off of coatings.
From Table 1, these alloys contain iron in various amounts from 0 to 1.8% by weight. From
Table 2, it is observed that all the alloys passed through
a large amount of undercooling in DTA experiments. The least amount of undercooling (T =
1650C) was observed in alloy NHFA-3 containing the largest amounts of boron, silicon and iron
while the largest amount of undercooling (T = 2910C) was recorded in alloy NHFA-1
containing no iron addition. Hence it could be safely said that the presence of iron in these alloys
played a crucial role in degree of undercooling of these alloys. This could be the basis of various
forms of solid-state precipitations observed in these types of alloys with iron additions.
3.2 Precipitations in the Ni (α) Solid Solution
Large segregations of silicon in the form of the silicide β Ni3Si were observed along the edges of
the Ni (α) solid solution (Fig.1b). The slow cooling rate gave rise to reversible process for
forming thermodynamically stable phases. The diffusion coefficients of the element diffusing in
the Ni (α) may be expressed as:
D = D0exp[-Q/kT]
where Q = activation energy in Cal.mol-1
T = temperature in K
According to Adda and Philibert [18], DSi = 1.5exp[-61700/kT] between 1100 and 13000C.
Hence, from the above relations, the diffusion coefficient of silicon (DSi) in nickel could be
estimated as 1.48 x 10-9cm2.s-1. The precipitates observed in the alloys (except NHFA-1 alloy)
were as a result of large segregations of silicon in the Ni (α) phase. This could be explained by
the solute exchange of silicon between the Ni (α) and the remaining liquid during cooling. This
resulted in the concentration profile of silicon being slightly higher in the Ni (α) solid solution as
shown in the micrograph of Fig. 1b. Similar observations have been reported in the nickel-rich
Ni-Si binary alloys [19]. Detailed microstructural observations showed that these precipitates
exhibited two characteristic morphologies (cuboids and spheroids) coherent with the Ni (α)
matrix as depicted in the micrograph of Fig. 4a. These morphologies have been explained by
Ricks et al [20] by the sign of the lattice misfit between the matrix and the precipitates and the
supersaturation of the diffusing species. Lebaili et al [15] reported the evolution of the cuboid
Vol.9, No.2 Scanning Electron Microscopy of Some Slowly Cooled Nickel–Based Hardfacing Alloys 141
precipitates into spherical precipitation and explained this by the growth of β Ni3Si along
preferential planes. This is in good agreement with the observations in this work. In fact the
occurrence of this type of precipitation has been extensively investigated in the case of γ (Ni3Al)
precipitates in the nickel matrix as it constitutes the strengthening mode for super alloys [20].
The presence of these precipitates has also enhanced the mechanical properties of these alloys as
evidenced by the dislocation loops around the precipitates (Fig. 4b) in the alloys under
investigation thereby increasing the strength of the alloys.
Figure 4. (a) Transmission electron micrographs (TEM) showing the cuboid and spheroid
precipitates within the Ni(α) matrix and (b) dislocations loops around the precipitates.
Other secondary boride precipitation observed in the nickel matrix made this matrix to be
particularly interesting to study (Figure 1b). From qualitative analysis, it was observed that these
needle-like precipitates within the volume of Ni (α) matrix were rich in boron, chromium and
tungsten. They are the boride of the type Cr2B. In fact from the nominal composition of this
142 J.A. Ajao Vol.9, No.2
alloy, carbon, boron and tungsten are rather not in appreciable quantities; hence they segregated
within the Ni (α) matrix. There was no evidence of the presence of chromium boride with higher
boron content such as CrB in all the alloys under investigation. During cooling, the solidification
of the M7C3 and MB phases from the liquid alloys occurred at sufficiently higher temperatures
(12700C to 12300C) than the temperature range in which the matrix components solidified
(9980C to 11400C). For most of the matrix components, the M7C3 and MB crystallized as the
primary and secondary phases so as to be dispersed in the matrix before the solidifying
temperatures of the matrix components were reached. Since the M7C3 and MB are both hard,
their nature, distribution, form, dimensions and crystallographic orientation relationship (mutual
or otherwise) are responsible for the wear resistance property of the alloys. It should be noted
that the degree of undercooling in these alloys was considerably reduced compared to Ni-B
binary [16] and Ni-B-Si ternary [21] alloys. This could be associated with the presence of
chromium which increased the liquidus temperature of the alloys.
3.3 Hardness of the alloys
The hardness values of the alloy are presented in Table 4.
Table 4. Average hardness values of the alloys under investigation.
Alloy HV10
NHFA-1 290
NHFA-2 335
NHFA-3 430
NHFA-4 415
NHFA-5 455
A plot of the hardness values of the alloys as a function of the iron contents is shown in Fig. 5.
From the figure, it can be seen that the hardness values increased as the iron contents increased.
Alloys NHFA-3 and NHFA-4 have close hardness values. This is probably due to the high
tungsten, boron and silicon contents in the alloys.
Vol.9, No.2 Scanning Electron Microscopy of Some Slowly Cooled Nickel–Based Hardfacing Alloys 143
-0,20,0 0,2 0,4 0,60,8 1,0 1,2 1,4 1,6 1,8
Iron contents wt%
Figure 5. Average hardness values of the alloys as a function of the iron contents.
On the other hand, alloy NHFA-1 has relatively low boron and silicon contents thereby
presenting low hardness value. The wide gap between the hardness values of alloys NHFA-2 and
NHFA-5 could be attributed to the low iron and tungsten contents in alloy NHFA-2. The high
hardness value of NHFA-5 could be associated with the high concentration of the hard carbide,
silicide and boride phases within the matrix due to the high contents of all the metallic and non-
metallic additions.
The morphology of some nickel-based hardfacing alloys has been studied by differential
thermal analysis (DTA), energy dispersive X-ray analysis (EDXA), X-ray diffraction
(XRD), scanning electron microscopy (SEM) and transmission electron microscopy
Three major primary phases (Ni (α), M7C3, and the π phase) were observed during slow
cooling in DTA depending on the nominal compositions of the alloys. Intense
segregations of the silicide phase Ni3Si were observed in the nickel matrix.
Large undercoolings in addition to intense solid state precipitations were observed in
alloys with iron additions. The precipitates are responsible for the high hardness values
and strength of the alloys.
The hardness values of the alloys were found to increase as the iron contents increased.
144 J.A. Ajao Vol.9, No.2
The author is indebted to Dr. Marianna Kemell for assistance in electron microscopy and to
Centre for Energy Research and Development, Obafemi Awolowo University, Ile-Ife, Nigeria
for granting him leave of absence during the preparation of this work.
[1] Lugscheider, E., Knotek, O. and Klohn, K., 1978, “Development of Ni-Cr-Si base Filler
Metals”, J. Weld. Res. Supp., Vol. 57 pp. 319 - 325.
[2] Gonz´alez, R., Cadenas, M., Fern´andez, R., Cortizo, J. L. and Rodr´ıguez, E., 2007, “Wear
behaviour of flame sprayed NiCrBSi coating remelted by flame or by laser”, Wear Vol. 262 pp.
[3] Knotek, O. and Lugscheider, E., 1976, “Brazing Filler Metals based on Reacting Ni-Cr-B-Si
Alloys”, Welding Research Supplement pp. 314 – 318.
[4] Steffens, H.D., Ferrari, F., Sturlese, S.,and Pawlowski, L., 1992, “Fundamentals of Thermal
Spraying Technology”, Mats. Eng., Vol. 3(2) pp. 183 - 190
[5] Knotek, O., Reimann, H. and Lohage, P., 1981, “Reactions between Ni-Cr-B-Si matrixes and
carbide additives in coating during fusion treatment”, Thin Solid Films Vol. 83 pp. 361 – 367.
[6] Knotek, O., Lugscheider, E. and Wichert, W., 1978, “On the structure and properties of wear
and corrosion resistant Ni-Cr-W-C-Si Alloys”, Thin Solid Films Vol. 53 pp. 303 – 312.
[7] Rosso, M and Bennani, A., 1998, “Studies of new applications of Ni-based powders for
hardfacing processes”, PM World Congress Thermal Spraying/Spray Forming, pp. 524-530.
[8] Knotek, O., Lugscheider, E. and Reimann, H., 1975, “Structure of Ni-rich Ni-Cr-B-Si
Coating Alloys”, J. Vac. Sci. Technol. Vol. 12(4) pp. 770 - 772.
[9] Gonz´alez, R., Garc´ıa, M. A., Pe˜nuelas, I., Cadenas, Roc´ıo, Fern´andez, Ma. Del.,
Hern´andez Battez, A., Felgueroso, D., 2007, “Microstructural study of NiCrBSi coatings
obtained by different processes”, Wear Vol. 263 pp. 619–624.
[10] Zhao, W-M., Wang, Y., Han, T., Wu K-Y. and Xue, J., 2005, “Corrosion mechanism of
NiCrBSi coatings deposited by HVOF”, Surf. Coat. Technol. Vol. 190 pp. 293 – 298
[11] Sidhu, T. S., Prakash, S. and Agrawal, R. D., 2006, “Hot corrosion and performance of
Nickel - based coatings”, Current Science Vol. 90(1) pp. 41 – 47.
[12] Otsubo, F., Era, H. and Kishitake, K., 2000, “Structure and phases in Nickel-base self-
fluxing alloy coating containing high chromium and boron”, J. Ther. Spray Technol. Vol. 9 pp.
107 – 113
[13] Lee, C. H. and Min, K. O., 2000, “Effects of heat treatment on the microstructure and
properties of HVOF – sprayed Ni – Cr – W – Mo – B alloy coatings”, Surf. Coat. Technol. Vol.
132 pp. 49 – 57
Vol.9, No.2 Scanning Electron Microscopy of Some Slowly Cooled Nickel–Based Hardfacing Alloys 145
[14] Wang, B. Q. and Luer, K., 1994, “The erosion – oxidation behaviour of HVOF Cr3C2
NiCr cermet coating”, Wear Vol. 174 pp. 177 – 185.
[15] Lebaili, S., Durand-Charre, M. and Hamar-Thibault, S., 1988, “The Metallurgical Structure
of as – solidified Ni-Cr-B-Si-C Hardfacing Alloys”, J. Mat. Sci. Vol. 23 pp. 3603 – 3611
[16] Ajao, J. and Hamar-Thibault, S., 1988,Influence of Additions on the Solidification
Behaviour of Ni-B Alloys- Crystallography of the Ni - Ni3B eutectic”, J. of Mat. Sci. Vol. 23,
pp. 1112-1125.
[17] Gladyshevski, E. I. and Borusevich L. K., 1963, “Ternary system of Cr-Ni-Si”. Zh.
Neorgan. Khim, Vol. 8 pp. 1915-1918
[18] Adda, Y. and Philibert J. La diffusion dans les solides, Presses Universitaires de France,
[19] Lebaili, S. and Hamar-Thibault, S., 1984, “Solid State Transformations During Cooling in
the Ni-rich Portion of the Ni-Si system”, Z. Metallkde Vol. 75 pp. 764-770.
[20] Ricks, R. A., Porter A. J. and Ecob R. C., 1983, “The growth of γ precipitates in nickel-
based superalloys”. Acta Metall., Vol. 31, pp. 43-53.
[21] Lebaili, S. and Hamar-Thibault, S., 1984, “Equilibres Liquide-solide dans le system Ni-B-Si
dans la region riche en nickel”, Acta Metall. Vol. 35(3) pp. 701-710.