Journal of Alloys and Compounds 469 (2009) 270–275
Effect of bonding temperature on microstructure development
during TLP bonding of a nickel base superalloy
M. Pouranvari ∗ , A. Ekrami, A.H. Kokabi
Department of Materials Science and Engineering, Sharif University of Technology, P.O. Box 11365-9466, Tehran, Iran
Received 3 November 2007; received in revised form 15 January 2008; accepted 18 January 2008
Available online 10 March 2008
Abstract
The effect of bonding temperature on microstructure of a transient liquid phase (TLP)-bonded GTD-111 nickel base superalloy, using a Ni–Si–B
interlayer, was investigated. At low bonding temperatures microstructure of the joint centerline is controlled by B diffusion. However, at high
bonding temperature effect of base metal alloying elements on the joint microstructure development was more pronounced. Effects of bonding
temperature on the formation of secondary precipitates in diffusion-affected zone and isothermal solidification time were also analyzed. It was found
that above a critical temperature, formation of borides in diffusion-affected zone is suppressed and the time required for isothermal solidification
completion is increased.
© 2008 Elsevier B.V. All rights reserved.
Keywords: High-temperature alloys; Microstructure; TLP bonding
1. Introduction
Transient liquid phase (TLP) bonding considered as a preferred repairing/joining process for nickel base superalloys due
to its ability to produce near-ideal joints [1].
A typical microstructure of TLP-bonded joint of a nickel
base precipitation hardened superalloy such as GTD-111 using a
boron containing interlayer, consists of three distinct microstructural zones, before completion of isothermal solidification [2]:
(i) Athermally solidified zone (ASZ), which usually consists
of eutectic microconstituents. This zone is formed due to
insufficient time for isothermal solidification completion.
Cooling is the main driving force for athermal solidification
(i.e. non-isothermal solidification).
(ii) Isothermally solidified zone (ISZ), which usually consists
of a solid solution phase. Compositional change induced by
interdiffusion between substrate and interlayer during holding at a constant bonding temperature is the driving force
for isothermal solidification. As a result of the absence of
∗
Corresponding author. Tel.: +98 912 4075960; fax: +98 21 66005717.
E-mail address: mpouranvari@yahoo.com (M. Pouranvari).
0925-8388/$ – see front matter © 2008 Elsevier B.V. All rights reserved.
doi:10.1016/j.jallcom.2008.01.101
solute rejection at the solid/liquid interface during isothermal solidification under equilibrium, formation of second
phase is basically prevented [3].
(iii) Diffusion-affected zone (DAZ), which consists of boride
precipitates due to B diffusion into the base metal (BM)
during TLP bonding.
Although some researchers [3–6] worked on TLP bonding of
␥ strengthened nickel base superalloys, but their works focused
on microstructure control of joint centerline via control of process parameter (e.g. bonding time, bonding temperature and joint
gap) to achieve a eutectic-free joint. There is a little emphasis
on effect of bonding variables on DAZ precipitates. The aim of
this research is to investigate and analyze the effects of bonding
temperature on microstructural features with particular emphasis on the DAZ precipitates and the time required for isothermal
solidification completion (tIS ) during TLP bonding of GTD-111
nickel base superalloy.
2. Experimental procedure
The chemical composition of the base metal, GTD-111 superalloy,
was Ni-13.5Cr-9.5Co-4.75Ti-3.3Al-3.8W-1.53Mo-2.7Ta-0.23Fe-0.09C-0.01B.
A commercial Ni-4.5Si-3.2B alloy (MBF30), in the form of an amorphous foil
with 25.4 ␮m thickness was used as the interlayer. The surfaces to be bonded
M. Pouranvari et al. / Journal of Alloys and Compounds 469 (2009) 270–275
271
Table 1
TLP bonding tests matrix
Bonding temperature (◦ C)
Bonding time
1100
1150
1180
30
30
30
45
45
60
75
60
120
were ground by using 600 grade SiC paper and cleaned in acetone before bonding. Bonding was carried out at 1100, 1150 and 1180 ◦ C for various bonding
times under a vacuum of approximately 10−4 Torr in a vacuum furnace. The
bonding schedule is given in Table 1.
Microstructures of joints were examined by optical microscopy and scanning electron microscopy (SEM) equipped with a beryllium window energy
dispersive spectrometer (EDS) system using INCA software. For microstructural
examinations, specimens were etched using two etchants. The Murakami etchant
(10 g KOH, 10 g K3 [Fe(CN)6 ], 100 ml H2 O) preferentially etches Cr-rich phases
and can therefore be used to reveal precipitates adjacent to the joint/base metal
interface. Molybdenum-acid etchant (0.5 g MoO3 , 50 ml HCl, 50 ml HNO3 ,
200 ml H2 O), which preferentially etches ␥ phase, was used to indicate ␥-␥
microstructure of the joints, in addition to joint centerline microstructure.
3. Results and discussion
3.1. Effect of bonding temperature on joint microstructure
Fig. 1 shows SEM image of joint made at 1100 ◦ C for
30 min. According to pervious work [2], microstructure of the
ASZ consists of a nickel-rich boride and a nickel-rich solid
solution, which form via a eutectic-type transformation and
microstructure of the DAZ consists of extensive blocky and
needlelike carbo-Cr-rich boride precipitates [2]. Morphology of
these precipitates is different from morphology of normal eutectic microstructure. This can be an evidence for the fact that these
precipitates were formed at the bonding temperature not during
cooling of bond to room temperature.
Fig. 1. Microstructure of bonds made at 1100 ◦ C for 30 min [2].
Fig. 2. Microstructure of bonds made at (a) 1150 ◦ C and (b) 1180 ◦ C for 30 min.
Fig. 2a shows a typical microstructure of joint made at
1150 ◦ C for 30 min illustrating some eutectic-type structure.
EDS analysis of the eutectic microconstituent showed that
microstructure constituent are similar to those of at 1100 ◦ C.
Fig. 2b shows microstructure of joint made at 1180 ◦ C for
30 min. Detail of joint centerline microstructure is shown in
Fig. 3 indicating that the joint centerline comprised of ␥-␥
eutectic and an intermetallic phase, which formed adjacent to
the non-equilibrium ␥-␥ eutectic. EDS analysis of the intermetallic phase showed that this phase is Cr-rich boride. Inset
of Fig. 3 shows microstructure of ISZ/ASZ interface indicating high volume fraction of gamma prime in gamma matrix.
␥ formed on cooling via solid state transformation, when temperature is lower than the ␥ solvus temperature. This can be
related to diffusion of more Ti and Al into the ISZ at 1180 ◦ C
in comparison to the lower bonding temperatures of 1100 and
1150 ◦ C.
One interesting feature of bonds made at 1180 ◦ C was the lack
of secondary precipitates, Cr-borides, in the DAZ (see Fig. 2b).
The lack of these precipitates was confirmed by Murukami
etchant. However, some liquated regions were observed in the
substrate region.
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M. Pouranvari et al. / Journal of Alloys and Compounds 469 (2009) 270–275
Fig. 3. Detail of joint’s centerline microstructure of bonds made at 1180 ◦ C for 30 min.
3.1.1. Boride precipitates in the substrate region
There are two explanations/mechanisms for boride formation
in the DAZ:
(a) Gale and Wallach [7] have suggested that contrary to the
predictions of currently available TLP models, a significant
solid state diffusion of B occurs in the BM before completion of the dissolution process. When concentration of B
in the BM exceeds from solubility limit of B at bonding
temperature, precipitation of secondary precipitates at the
joint/BM interface is expected.
(b) The second mechanism is based on the diffusion of B into
the BM during isothermal solidification. Isothermal solidification process progresses by B diffusion from liquid into
the BM. Diffusion of boron from liquid into the BM during isothermal solidification and presence of Cr, which is
a strong boride former, can explain formation of Cr-rich
borides. According to this mechanism, boron concentration
profile across the joint and the resultant microstructure can
be schematically shown as Fig. 4a. The boron concentration in the ISZ is above the boron solubility limit in the
base metal. Local increase in boron content in the DAZ is
due to the boride precipitation in this region. The amount
of B, which can diffuse through the ISZ into the BM, is
governed by B solubility at the liquid/ISZ interface (C␥L ).
If boron content in the ISZ be greater than boron solubility
in the BM (CSBM ), boride precipitation in the BM would be
expected.
Although bonding temperature of 1180 ◦ C is still lower than
the borides solvus temperature, no Cr-borides were observed in
the DAZ. This phenomenon can be explained as follows:
(i) According to the Gale and Wallach’s mechanism [7] if the
dissolution process is completed and the solid and liquid
phases are equilibrated (with their respective solidus and liquidus compositions) before significant solid state diffusion
has taken place, then the behavior of the Ni–Si–B interlayer
should be similar to that assumed in the conventional TLP
models. Nishimoto et al. [8] applied the Nernst–Brunner
theory for modeling the dissolution stage during TLP bonding:
kt
(1)
C = Csat 1 − exp −
W0
where C: MPD concentration at the time t; Csat : MPD
concentration in the liquid at equilibrium; k: dissolution
rate constant, W0 : interlayer thickness; t: holding time.
According to Nernst–Brunner theory, time required for the
base metal dissolution completion reduced with increase
in dissolution rate constant. Dissolution rate constant (k)
increases with increasing bonding temperature according
to the temperature-dependent Arrhenius expression. Consequently, an increase in bonding temperature should result
in shorter dissolution times. Therefore, time available for
diffusion of B out of liquid interlayer into the base metal is
reduced. This can be contributed to the suppression of the
boride formation in DAZ.
(ii) According to the second mechanism, if boron concentration in the ISZ be less than the boron solubility in the
BM (CSBM ) boride precipitation in the BM would not be
expected. In this condition, the formation of borides is suppressed because the boron content in the base metal material
cannot exceed the solubility limit. On the one hand, presence of Cr, W and Mo, strong boride forming elements,
M. Pouranvari et al. / Journal of Alloys and Compounds 469 (2009) 270–275
273
elements (6.5 wt% Cr, 6 wt% W and 0.6 wt% Mo), is about
1289 ◦ C [9]. Therefore, considering substantial content of
these elements (13.5 wt% Cr, 3.8 wt% W and 1.53 wt% Mo)
in GTD-111, it can be deduced that solubility of B in the
BM increases with increasing the bonding temperature from
1100 to 1180 ◦ C. Therefore, considering these two factors
(i.e. decreasing of B concentration in ISZ and increasing
of B solubility in BM with increasing the bonding temperature from 1100 to 1180 ◦ C) it seems that boron solid
solubility of the ISZ decreases to the values lower than solid
solubility of boron in the BM at bonding temperature of
1180 ◦ C, as schematically shown in Fig. 4b. Consequently,
according to the second mechanism a joint interface free
from boride precipitates can be expected. A similar scenario was proposed by Schnell [9] and Schnell and Graf
[10] who reported a reduction in volume fraction of borides
at high bonding temperature during TLP bonding of CMSX4 using a Ni–Cr–Co–Al–Ta–B filler alloy. They attributed
this phenomenon to the fact that above a critical temperature, the boron solubility in the isothermally solidifying
layer is reduced down to the boron solubility of the parent
material.
Fig. 4. Boron concentration profile and the resultant joint microstructure for
bonding temperature of (a) 1100 ◦ C and (b) 1180 ◦ C.
reduces B solubility in the nickel base alloys. Therefore,
due to more diffusion of these elements from BM into the
ISZ at higher bonding temperature, B solubility of in ISZ
is reduced (i.e. C␥L is reduced). Maximum solid solubility
of B in Pure Ni is at temperature about 1094 ◦ C. W and Mo
shift the boride solvus temperature upwards [9]. Therefore,
on the other hand, since the GTD-111 has substantial Cr,
W and Mo, the temperature of maximum solid solubility
of B in the BM is higher than that in the pure nickel substrate. Through modeling it has been shown that maximum
solubility temperature of boron in CMSX4 nickel base single crystal superalloy, which has substantial boride forming
3.1.2. Solidification behavior in the joint centerline
Solidification behavior of ASZ is governed by formation of
primary ␥ solid solution and subsequent solute partitioning.
Microstructure of ASZ in bonds made at 1100 and 1150 ◦ C
is governed by segregation of MPD elements during nonequilibrium solidification. Very low solubility of B in Ni and
partition coefficient of B in Ni [11] lead to rejection of B into
the adjacent melt shifting composition of melt towards the eutectic composition; thus, with progressing of solidification, binary
eutectic of ␥-solid solution and nickel boride is formed. The EDS
analysis of ASZ showed no silicide phase in this zone. Indeed,
low solubility and low partition coefficient of B in Ni base
substrates resulted in formation of intermetallic phases within
eutectic structure of ASZ. Therefore, it can be deduced that joint
microstructure and isothermal solidification rate at 1100 and
1150 ◦ C is dominated by diffusion of B.
According to the microstructural investigation, at bonding
temperature of 1180 ◦ C, the ASZ microstructure is controlled
by segregation of BM alloying elements particularly Ti and Cr.
Indeed, more dissolution of BM alloying element and more diffusion of Ti and Cr from BM during isothermal solidification at
bonding temperature of 1180 ◦ C in comparison to lower bonding temperatures of 1100 and 1150 ◦ C govern the microstructure
development in the ASZ. It has been reported that ␥-␥ eutectic
formed due to the segregation of some alloying elements particularly Ti during non-equilibrium solidification of nickel base
superalloy [12]. Cr, Mo, and W have low concentrations in ␥
phase, the main phase in ␥-␥ eutectic; therefore have low concentrations in ␥-␥ eutectic [13]. Therefore, residual liquid is
enriched with these boride forming elements. It has been also
reported that Cr and Mo exhibit positive segregation into the
residual liquid phase during ␥-␥ eutectic transformation in a
nickel base superalloy [14]. Therefore, considering enrichment
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M. Pouranvari et al. / Journal of Alloys and Compounds 469 (2009) 270–275
Fig. 5. Optical micrograph showing fracture path in bonds made at 1180 ◦ C for
30 min.
of liquid with boride forming elements (e.g. Cr, Mo and W)
and presence of boron in the liquid phase, formation of Cr-rich
borides in the ASZ can be explained.
Tie up of Ti, as a key alloying element in ␥ (Ni3 Al, Ti);
in ␥-␥ eutectic can significantly affect high-temperature performance of the joint. More importantly, as can be seen in Fig. 5
presence of non-isothermally solidified products in the joint
centerline can be a preferential source of failure and crack propagation occurred through the ␥-␥ eutectic and Cr-rich boride.
Despite the fact that there are significant ␥ precipitates in ISZ
microstructure, the shear strength of this joint was very low
(about 30% of BM shear strength). Indeed, the joint strength controls by nature and size of ASZ, when isothermal solidification
is not accomplished completely.
3.2. Effect of bonding temperature on the time required for
isothermal solidification completion
To investigate the influence of bonding temperature on tIS ,
bonding was carried out for different bonding times at 1100,
1150 and 1180 ◦ C. The average ASZ widths were measured
using SEM micrographs and plotted against bonding time for
various bonding temperature (Fig. 6). As can be seen, the ASZ
width decreases with increasing bonding time at a given bonding temperature. Isothermal solidification, which prevents the
formation of centerline eutectic constituent, completes at the
bonding time of 75 min at 1100 ◦ C. Increase in bonding temperature is expected to significantly reduce the required time
to attain a eutectic-free joint due to the increase in MPD diffusion coefficient. Sakamoto et al. [15] proposed a linear relation
between maximum clearance free from eutectic and a brazing
parameter, (T + 20 log t)1/2 , where T is bonding temperature and
t is bonding time. Therefore, it was suggested that the increasing
bonding temperature results in shorter tIS .
Although increasing the bonding temperature to 1150 ◦ C led
to decreasing of tIS to 45 min (see inset of Fig. 6), a eutecticfree joint was not attained even after 120 min holding time at
1180 ◦ C (see inset of Fig. 6). Therefore, isothermal solidification
behavior at 1180 ◦ C is abnormal with regard to conventional
Fig. 6. Variation of eutectic structure size with bonding time at various bonding
temperature.
expectation. This is in agreement with the previous works on
TLP bonding of In738LC using Ni–Cr–B interlayer by Idowu
et al. [5] and Wikstrom et al. [6] who reported that isothermal
solidification rate decreases at high temperatures.
Isothermal solidification requires diffusion of MPD element
out of liquid phase into the base metal until liquidus temperature
of the liquid phase reaches the bonding temperature. When liquidus temperature of liquid phase reaches bonding temperature
isothermal solidification starts. Isothermal solidification accomplished completely when liquidus temperature of residual liquid
at the center of the joint reaches the bonding temperature. In this
situation MPD concentration at the center of the joint is reduced
to the solid solubility of MPD at the bonding temperature [5,6].
Therefore, tIS depends on the MPD elements flux from liquid
phase into the BM and phase relationships in BM/interlayer system. In binary TLP system, a critical bonding temperature is
also expected due to interplay between increase in diffusivity of
MPD elements and increase in maximum widening of the liquid [16]. However, in multicomponents systems such as present
system, GTD-111/MBF30, this cannot be the sole explanation.
The observed behavior can be explained as follow:
(i) As mentioned previously, B controls isothermal solidification process at 1100 and 1150 ◦ C, while at bonding
temperature of 1180 ◦ C isothermal solidification is controlled by Ti. Diffusion coefficient of Ti in Ni base alloys
is much lower than that of B in Ni. Ti is a substitutional
solute and B is interstitial solute in Ni.
(ii) The solidification partition coefficient of Ti in nickel is
less than unity [17]; therefore, an increase in its concentration in the liquid would result in a depression of residual
liquid solidification temperature. Isothermal solidification
accomplished completely when liquidus temperature of
residual liquid at the joint centerline reaches the bonding
temperature, as a result of diffusion of MPD elements from
liquid to the adjacent solid. Therefore, considering low diffusion coefficient of Ti in Ni and high difference between
bonding temperature and liquidus temperature of resid-
M. Pouranvari et al. / Journal of Alloys and Compounds 469 (2009) 270–275
ual liquid, lower isothermal solidification rate at 1180 ◦ C
compared to the isothermal solidification rate at 1100 and
1150 ◦ C can be explained.
(iii) In addition, boride precipitation in the DAZ can affect the
diffusion of boron during isothermal solidification. Fraction of boron consumed by boride formation and thus taken
away from the matrix phase enables more boron to enter
within the base material, which has an accelerating effect
on the effective boron diffusion [18]. As seen in the previous section, at 1180 ◦ C no boride phase was observed
in the BM, hence effect of boride formation on the boron
flux is negligible. This can also attribute to the reduction
of isothermal solidification rate at 1180 ◦ C.
Isothermal solidification kinetics during TLP bonding
depends not only on diffusivity of MPD elements but also on
the MPD solubility in the base metal. Indeed, in multicomponents system such as nickel base superalloy interplay between
these factors is the controlling factor for the required time to
attain a eutectic-free TLP joint. Therefore, developing of multicomponents TLP models is necessary to establish a better
understanding of isothermal solidification during TLP bonding
of complex alloy systems.
According to the result of this study, to achieve a proper
bond microstructure bonding temperature should be optimized.
Increasing bonding temperature higher than a critical temperature has an opposite effects on the tIS and extent of secondary
precipitates formation in the joint interface. The condition in
which isothermal solidification has not been accomplished completely, eutectic structure which has usually the highest hardness
in the bond region and forms an interlinked network is the preferential failure source. Secondary precipitates in the DAZ have less
detrimental effect on joint strength. However, high chromium
content of DAZ precipitates can lead to a significant depletion
of chromium around this region of the BM, which may result
in a decrease in the corrosion resistance. Therefore, considering mechanical properties of the joint, bonding at temperatures
higher than the critical temperature, at which substantial diffusion of BM alloying elements (particularly for Ti) occurs,
should be avoided. In order to eliminate Cr-rich precipitates
from DAZ and restore corrosion resistance of this region it is
better to design a proper post-bond heat treatment after isothermal solidification completion rather than to use high bonding
temperature.
4. Conclusions
From this research the following conclusions can be drawn:
(1) Bonding temperature affect significantly athermally solidified zone microstructure; at lower temperature ASZ
microstructure is governed by B diffusion, while at high
275
bonding temperature ASZ microstructure is significantly
affected by dissolution and diffusion of BM alloying element into the joint region, particularly Ti and Cr.
(2) DAZ precipitate formation is significantly affected by bonding temperature. At bonding temperatures of 1100 and
1150 ◦ C extensive Cr-rich carbo-boride was observed in
DAZ, while at bonding temperature of 1180 ◦ C joint/BM
interface was almost precipitate-free.
(3) There is a critical bonding temperature to minimize the
required time for isothermal solidification and, hence, to
attain a eutectic-free TLP joint. Decreasing of isothermal
solidification rate at bonding temperature of 1180 ◦ C can
be attributed to the following factors, in addition to increasing joint width: (a) enrichment of residual liquid with some
BM alloying element particularly Ti and Cr, during dissolution and isothermal solidification, (b) reduction of boron
flux into the BM at this bonding temperature due to lack of
boride precipitates in the BM.
Acknowledgment
The authors would like to acknowledge Sharif University of
Technology for financial supporting of this research.
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