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Diffusion in the Interface Region of Ti/TiAl-Nb Bonding
Lembit Kommel
Department of Materials Engineering, Tallinn University of Technology
Ehitajate tee 5, 19086 Tallinn, Estonia, E-mail:kommel@edu.ttu.ee
Keywords: Diffusion bonding, Titanium/aluminum intermetallides, Interface region.
Abstract. Diffusion in the interface region of advanced metals, light weight heatproof quality
titanium alloy and titanium/aluminum intermetallide alloy, was investigated. We studied the
diffusion of aluminum from intermetallide to titanium alloy and change of other chemical elements
and microhardness in diffusion region; it was formed in solid titanium alloy. The studies of
microstructure showed that the interface region included a transition zone in the initial solid Ti-
alloy and the second molten TiAl-Nb intermetallic substrate. The diffusion interface region of
diffusion bonding on optical light microscope was 45-60 m in width. By this the titanium content
decrease and aluminum content increase from titanium alloy surface up to 120-150 m in depth of
solid titanium alloy. As result of diffusion the intermetallide Ti
3
Al thin layer was formed in the
transition zone in Ti-alloy substrate, and also the diffusion microporous were formed in interface
region.
Introduction
Light weight Ti-alloys and intermetallide TiAl-Nb-alloys are widely used as constructional
materials in aerospace [1], in turbo-jet [2] and in internal combustion engines [3, 4] application.
Titanium aluminides [5-8] (depending on chemical composition, microstructure, crystallite size, and
processing mode) one characteristics are: low density, high temperature specific strength, good
oxidation and burn resistance, excellent creep and relaxation properties. It means that this material
single crystal [9] has good mechanical performance at elevated temperatures, resistance to oxidation
and low specific density.
The titanium aluminides [1], for example, can be used in thermal protection systems (TPS) for next-
generation low orbit and re-entry space vehicles. Unfortunately this material manufacturing is very
difficult. However, the major obstacle to TiAl-intermetallide alloys applications is the low ductility
at room temperature [1-5] and during hot forging [10]. It is mean that the TiAl-intermetallide
structure alloys are difficult to process mechanically. For example, in contrast to titanium alloys, the
productivity of the turning processing of the details from intermetallide alloys is reduced up to ten
times. Thus, the good machining results are received at polishing and electrical erosion processing
of intermetallides.
So to weld the heat proof TiAl-Nb intermetallide alloy to Ti-alloy to form compound structure is
necessary [3, 4, 11]. It can reduce the manufacturing time and increase the work temperature of
component. Therefore, study of the different joining methods [12-15], formed microstructure,
microhardness and mechanical properties in interface region are important.
To ensure good plastic properties of TiAl-intermetallide alloys at room temperature, different
manufacturing technologies have been developed. For example, the directionally solidified ingots
[8] of TiAl-intermetallide have aligned lamellar microstructure. This intermetallide structure
exhibits a high yield stress with reasonably large tensile elongation at room temperature. Using the
HIP-process [5] at 900°C by pressure at 200 MPa, the diffusion bonding of the Ti-alloy to the -
intermetallide was formed. The tensile and creep properties were investigated at temperatures
between the room temperature and 600°C. Creep testing showed that most of the creep elongation
occurs in the Ti-alloy, but failure is initiated in the joint bond line. Creep causes degradation and
pore formation in this line. Interlining of these pores creates a crack which grows slowly until the
fracture toughness of the -TiAl is exceeded and the crack starts to propagate in the metal
terminating creep life. Results of investigation [12] show that during rapid solidification of the melt
under laser pulse-pressure, the periodic ablation structure of a metal is formed. For samples of TiAl-
intermetallide [14], obtained as a result of impulse processing during super cooling, the wave type
distribution of the strain fields and microhardness also reflects the specifics of interaction and multi-
level hierarchy of defects caused by additional energy take-off. The diffusion in the Ti-Al system as
on atomic mechanism of diffusion in the compounds is analyzed in detail, and methods of diffusion
calculations under different mechanisms are investigated in [15]. The high-temperature stability of
the lamellar structure and creep resistance are determined by diffusion rates in the phases and along
the interfaces. These experimental methods, used in modern diffusion mechanism, are briefly
described.
To achieve the complex of properties of the TiAl-intermetallide, the impulse fused-forging
modeling (IFFM) process [12] was invented. According to this process, the smelt of the metal can
be received by means of electro-arc melting, remove the smelt into the die and following impact
pulse action on it during solidification under high pressure. This process for dual engine valves with
TiAl-intermetallide disc on Ti-alloy steam manufacturing was used [4]. It was shown that the
summarized tangential thermal stresses and pressure from the valve’s head body depends on the
temperature level and physical properties of interface metals. The coefficients of heat transfer and
thermal expansion and the heat capacity of TiAl-intermetallide and Ti-alloys are different. The
thermo-physical parameters of these materials increase identical by temperature increases up to
650-700°C.
Taking into account these results of the previous studies, materials for specimens manufacturing
of the TiAl-Nb intermetallide and heat proof titanium alloy were chosen. For specimens
manufacturing we used the IFFM-process [12, 14]. The region of the interface, the microstructure
and microhardness, chemical content and phases, microporous and cracks forming were
investigated and analyzed.
Experimental
The IFFM-process technology is described in details in [12]. By this technique the intermetallide
plank in upper camber can be melted by electron beam method under inert gas argon pressure. The
contra part of solid Ti-alloy situated in bottom vacuum camber. These cambers were coined by
pipeline and situated on distance of one meter. When the intermetallide plank was melted in upper
camber the pressure was increased up to maximum pressure. As result of gas pressure difference in
the upper and bottom cambers the melted intermetallide was moved by speed up to 23.5 m/s into
water-cooled stamp with Ti-alloy substrate sample.
For the solid blank material, the heat-resistant (-titanium alloy (Ti-6.5Al-3.7Mo-3.4Zr-0.4Si)
and for the remelting material the TiAl-Nb-intermetallide (Ti-37Al-6Nb) blank were used.
We analyzed the microstructure of these materials and of diffusion bonding in the interface region,
resulting from the IFFM-process by means of the optical microscope (OM) Nickon CX, scanning
electron microscopes (SEM) JOEL JSM-840 A and Gemini, LEO, Supra 35. To analyze chemical
and phase structure in the interface region, we used the energy-dispersive X-ray microanalyses
(JOEL 840/LINK ANALYTIC 95/10 and Gemini, LEO, Supra 35) and X-ray diffractometer
(D5005, Bruker AXS). The micro indentation method for microhardness of phases of diffusion
bonding and of joined metals was used. By help of a microhardness tester Micromet-2001, Buehler
at a load of 50 gr during the testing time of 12 second the microhardness was measured. Test results
were computed and used for the analysis of diffusion behavior in bonding region.
Results and discussion
Microstructural investigation. Different microstructures in the optical and electronic pictures (Fig.
1, a, b) of diffusion bonding with the interface region and microstructures near bonding (Fig. 2, a, b)
are shown. The right side of the photo shows the -phase (after polymorphous transformation) of
Ti-alloy. The left side presents the TiAl-intermetallide. The middle part of the photo shows the
diffusion bonding region. As a starting condition, Ti-alloy has a globular (+)-phase structure.
a) b)
Fig. 1. The optical (OM) and electronic (SEM) pictures of microstructure of bonded Ti/TiAl-Nb
interface region are shown.
a) b)
Fig.2. SEM pictures of intermetallide (a) and of titanium alloy (b) transformed microstructures near
diffusion bonding.
As a result of heating on short time over the polymorphous transformation temperature (T
c
), the
microstructure of Ti-alloy was changed from (+) to -phase. The polymorphous transformation
of Ti-alloy microstructure occurred in the depth up to 30 mm along the specimen. By this the zone
of polymorphous transformation from - to -phase has duration of about 200 m in length. The
microstructure of TiAl-Nb intermetallide (Fig. 1, a, left side) is a congestion of
2
-, - and (
2
)-
scale of phases. These structural formations in intermetallide disc are oriented at different corners.
The maximal phase’s sizes are up to 100 m. Thicknesses of alternate plates of the phases are less
than one micrometer. These plates are a basis of the phase of intermetallide alloy. These phase’s
transformations of Ti-alloy near diffusion bonding depend on the temperature of intermetallide
smelt and the heat transfer coefficient between of the intermetallide smelt and Ti-alloy at the
moment of impact and, during solidification and cooling. The cooling rate during processing had a
major effect on the structure forming of TiAl-Nb intermetallide.
The combination of impact pressure and speed of solidification of the melt results in a metal of fine-
grained structure. During the fusion process, the melt is in contact with its own solid phase of blank
for remelting only. The thermal balance of the system as a first approximation can be written as Q
1
= Q
2
+ Q
3
, where Q
1
– heat addition from under-cooled intermetallic smelt, Q
2
– heat receiver at Ti-
alloy steam and Q
3
– heat extraction to water cooled-copper die. The interface quality or additive
property depends on the temperature and impact pressure level on the contact zone between smelt
and solid phases. By this the structure of the intermetallide depends on the impact pressure and
temperature of undercooled [15] smelt. These parameters are in a synergistically bound state. The
optical photo (Fig.1, a, left), dark and gray shades represent the two phases. This fact shows that
2
-
phase indicates the occurrence of rapid solidification at temperature about 1180°C, which does not
conform to the well known phase diagrams for these systems.
X-ray investigation results. The X-ray investigation show, that region of intermetallide has - and
2
-phases, accordingly. Studies by help of the energy-dispersive X-ray microanalyser of the
diffusion bonding region showed that the chemical content of major metals Ti and Al changed from
Ti-alloy condition up to the TiAl-Nb-intermetallide state. X-ray analyses were conducted on two
sides of the diffusion bonding. The intermetallic part of bonded metals exposes two major sets of
Bragg reflections. Each set corresponds to TiAl or -phase and Ti
3
Al or
2
-phase, respectively. The
ratio of the width of - (004/002) and
2
- (400/200) reflections is within the limits 1.30-1.38 of the
intermetallide alloy. This reflects the fact that broadening of Bragg reflections is a result of the
microscopic stresses caused by the impact pressure and diffusion mechanism of solid-state phase
transformations.
Fig. 3. X-ray investigation of intermetallide in diffusion bonding.
Indexed by 1997 JCPDS X-ray pattern of IFFM processed Ti-46Al-4Nb (wt. %) intermetallide:
1-TiAl, 2-Ti3Al, 3-Ti, 4-Al2Ti, 5-Nb, 6-Al3Nb, 7-Al, 8-Al3Ti, 9-Cu9Al4, 10-Al5Ti3.
The chemical content measure points and elements changes of the bonding are shown in (Fig. 4, a,
b). These results showed that changes in the titanium content take place from the distance 120 m
inside Ti-alloy from 86 wt. % down to 63.8 and 58.9 wt. % (
2
and phases, accordingly) on the
initial surface and intermetallide only. The titanium content did not change the chemical content of
intermetallide near diffusion bonding to Ti-alloy. Aluminium content increases from 6.5 wt. % (in
Ti-alloy) up to 37.5 and 31.7 wt. % (- and
2
- phases, accordingly) in intermetallide. Molybdenum
content in Ti-alloy decreased from 4.5 and 3.6 wt. % (in
2
- and - phases, accordingly) up to zero
in middle part of bonding. By this, niobium was diffused in Ti-alloy on the distance up to 40-50
m. The median chemical content (Fig. 4, up to the vertical line, the surface of the solid Ti-plank in
the starting condition) was: Ti60 wt. %, Al36 wt. % and Nb4 wt. %, accordingly. The -phase
element composition was: Ti58.9 wt. %, Al37.5 wt. % and Nb3.6 wt. %, accordingly. The
2
-phase
element composition was: Ti63.8 wt. %, Al31.7 wt. % and Nb4.5 wt. %, accordingly. The chemical
content of Ti and Al has an abrupt change in the vertical line. Ti content increases (up to 69 wt. %)
and Al content decreases (to 26 wt. %) in the distance of 1 m at the solid surface of titanium alloy.
Lin (Cps)
0
1
2
3
4
5
6
7
8
9
10
11
12
13
14
15
2-Theta - Scale
21 30 40 50 60 70 80 90
4
2, 9
1, 2, 3, 5, 7
1, 7, 10
1, 9
6
1, 2, 6, 7
1, 2, 9
1, 2, 6, 7
5, 10
6, 8, 10
3, 5, 7, 8
2, 3, 5
6, 8, 10
6, 8
6, 8
6 6
6
2
2, 9
2
2
2
2
2
2
2, 10
3
9, 10
9
9
9
10
10
10
10
a) b)
Fig. 4. The SEM investigation of chemical elements distribution in the diffusion bonding region: a –
chemical elements testing by line and test points of α
2
and γ- phases, test points 1, 2, and 3 on
different distance from interface; b – characters per second (cps) change of chemical
elements of spectra during SEM testing.
Fig.5. The calculated distribution of chemical elements (in wt. %) in diffusion bonding region of
TiAl-Nb intermetallide (left side) to solid Ti-alloy substrate (right side), accordingly.
In the distance of about 3 m, Ti content increases up to 73 wt. % and Al content to 20 wt. %,
accordingly. After this (by visual region width in this part about 62 m, Fig. 2) Ti and Al content
changes have a linear character up to 100-110 m the deepness of Ti-alloy. After this vertical
interface line, the content of alloying elements (3.7Mo, 3.4Zr and 0.4Si) of Ti-alloy and of TiAl-
intermetallide (4Nb) in total increase. The distance up to middle part of the diffusion region,
followed in the distance of 50-60 m from the decrease in the vertical line to about 7.5 wt. %, as in
Ti-alloy.
Microhardness testing. The results of measuring microhardness (pressure load 50 gr and
indentation time 12 s) are shown in Fig. 6. Microhardness of - and
2
-phases of TiAl-Nb-
intermetallide was studied in [2, 3]. According to this investigation (Fig. 6. curve - IFFM) the
microhardness of the diffusion bonding region was higher than ten year ago. By this, the chemical
content of elements has not changed.
Fig. 6. The microhardness distribution in diffusion bonding region of TiAl-Nb intermetallide (left
side) to solid Ti-alloy substrate (right side) immediately after IFFM process (curve IFFM),
microhardness after edging time at ten years (curve Ti) and microhardness of phases TiAl and Ti
3
Al
of intermetallide (left side), accordingly.
The relaxation decreases in microhardness in the diffusion region indicates that the broadening of
Bragg reflections is a result of the microscopic stresses caused by pressure and solid-state phase
transformation. The formation of new phases in the diffusion region is the result of atom diffusion.
When Ti, Al and Nb atoms diffuse to a certain extent near the interface surface of Ti-alloy part, the
chemical reaction occurs, and a new compound with yours microhardness is formed. When the
content of Ti is 80-85 wt. % and Al content decrease to 15-10 wt. % and the Nb content decreases
to 0.5-1 wt. % by the alloying elements (Mo, Zr, Si) of Ti-alloy concentration on the level 6-6.5 wt.
%, microhardness will increase up to the maximal value (Fig. 6, curve - IFFM). During the edging
time of ten years, microhardness will decrease more than two times. This decrease of microhardness
from 1050HV0.05 to 500HV0.05 is a result of stresses relaxation in the diffusion region of the
bonded alloys.
Diffusion defects analyze. As result of edging time of ten years the diffusion microvoids and
microcracks in the bonding were formed (Fig. 7). The micro defects were formed in result of
intrinsic stresses, which were formed during diffusion influenced by rapid solidification of
intermetallide smelt by IFFM process. The microvoids (Fig. 7, b and c) can be formed immediately
during IFFM as result of rapid diffusion of elements (Fig. 4). The microvoids do not formed in
interface (Fig. 1, a, b) where the temperature gradient is lower. Defects forming in the interface
region depend on physical properties of bonded materials also.
a) b)
c) d)
Fig. 7. Defects analyze affected by cooling rate on diffusion porosity formation and effect of micro
stresses on intergranular cracking in diffusion bonding region in (a), diffusion microvoids in (b),
microvoid in (c) and microcrack in (d) are shown.
Summary
1. TiAl-Nb intermetallic smelt was successfully connected to solid Ti-alloy by the IFFM technique.
2. The interface region of TiAl-Nb/Ti diffusion bonding includes a transition region in the Ti-alloy.
3. The TiAl-Nb intermetallic region includes the intermetallic compounds TiAl and Ti
3
Al, and new
by formed compounds such as Al
2
Ti, Al
5
Ti
3
, and Al
3
Nb.
4. The microhardness of diffusion bonding after a ten-year edging was only 500-520HV0.05.
5. The microhardness of - and
2
– phases were 275HV0.05 and 320HV0.05, accordingly.
6. Maximum microhardness of 1020HV0.05 was measured after the IFFM process immediately in
the region of diffusion bonding by the chemical content of Ti-80, Al-15, Mo-3.5, Zr- 3.2, Si-0.3 and
Nb-1 (in wt. %), accordingly.
7. Controlling microhardness and chemical compounds can reduce the width of Ti-alloy.
8. This dissimilar metals joint technology will be used in the aerospace or automobile industries to
manufacturing high temperature structural components from heavy duty advanced engineering
materials.
Acknowledgements: This work was supported by the Estonian Science Foundation Grant G-5878.
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