Abstract– Solid-state direct diffusion bonding of commercially pure Titanium (Cp-Ti) and precipitation hardening stainless steel (PHSS) has been carried out in the temperature range of 800°C to 1000°C with an interval of 50°C for 3.6 ks under 3.5 Mpa uniaxial load in (4 to 6)×10–3 Pa vacuum. The effects of temperatures have been investigated with reference to bond strength. The examination revealed that the transition joints achieve tensile strength 108% and shear strength 87.6% higher compared to pure titanium. And the ductility of the joint has been found to be 12.8% when processed at 950°C. The joint structure and intermetallic phases like Fe2Ti, FeTi, Cr2Ti, λ ( solid solution of Fe2Ti and Cr2Ti ), χ ( Fe17Cr7Ti5 ), α-Fe, α-Ti and β-Ti were predicted by optical microscope, SEM-BSE, EPMA and subsequently the presence of intermetallic phases including σ phase (compound of Fe and Cr ) were confirmed by X-ray diffraction (XRD) technique. EPMA also revealed substantial diffusion of Fe, Cr and Ni into the titanium base-metal and titanium traveled a less distance towards PHSS side. Bond strength and hardness of the interface layers were measured by tensile testing and micro-hardness testing respectively. The activation energy Q and growth velocity k of the reacting layers in the diffusion bonded joints were also calculated. The bond strength has been found to be satisfactory at 950°C. Bond strength drops as the volume fraction of intermetallics increased with the rise in bonding temperature (1000°C). Index Terms– Diffusion bonding, Intermetallics, SEM-BSE, EPMA, X-ray diffraction.
I. INTRODUCTION
In high-Tech. engineering application, however suitable bonding techniques are required to fabricate a composite structure by using multi components alloys like steels, Ti, Ti-alloys, super alloys etc. [1]. Conventional bonding of dissimilar material is much critical due to various metallurgical heterogeneities like thermal expansion, mismatches, large differences in melting points, development of residual stress and formation of brittle intermetallic phases. After much exploration of the processes, it has been suggested that
solid state joining process can give the better solution and in this respect, diffusion bonding is convenient with minimum macroscopic deformation and reduction of mechanical properties [2]. In solid state diffusion bonding, two materials are brought in close contact at elevated temperatures under moderate pressure. However, diffusion takes a dominant role with extent of diffusion zone and control of diffusion reaction, under the influence of parameters like temperature and pressure. Owing to low processing temperature, it is possible to eliminate the problem generated by fusion welding.
The bonding progresses in various stages; first stage, material contact areas of the mating surfaces as well as joint areas. In the second stage, diffusion of grain boundary predominates. This stage, eliminate pores and finally ensue the grain boundary arrangements. In third stage, volume diffusion dominates and joining process is completed. Titanium and precipitation hardening stainless steel have the wide applications in nuclear chemical, aircraft, naval industries due to their excellent mechanical properties and excellent corrosion resistance [3]. These two materials can be joined by diffusion bonding process [4]. According to the literature [5-10], at room temperature Ti and Fe have limited solubility in each other and formed intermetallic phases along with terminal solid solution in the binary phase diagram. Direct bonding between Ti and Stainless Steel with varying temperature promotes the formation of intermetalics like σ, Fe2Ti, Cr2Ti, Fe2TiO4, FeTi, λ and χ phases [5-10]. In present study, Solid-state direct diffusion bonding of commercially pure Titanium and precipitation hardening stainless steel (PHSS) has been carried out in the temperature range of 800°C to 1000°C and with an interval of 50°C for 3.6 ks. Under 3.5 Mpa uniaxial load in ( 4 to 6 ) ×10–3 Pa vacuum. The aim of the present investigation is revealing the different reaction layers formed near the interface and influence of bonding temperatures on the structural changes of the diffusion zones and mechanical properties of the bonded sample.
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II. EXPERIMENTAL
The commercially pure titanium (Cp-Ti) and PHSS, used in the present investigation were in the form of 25mm diameter rods. After hot rolling, Cp-Ti was annealed at 700°C and subsequently PHSS was solutionized at 1050°C and quenched, followed by tempering at 450°C for 14.4 ks. The heat treatment of PHSS has been carried out to produce precipitation of intermetallics compound that can cause precipitation hardening in stainless steel. The chemical compositions and mechanical properties after afore said heat treatment of the base metals are presented in Table 1 and Table 2. PHSS of these chemical compositions regarded as martensitic PHSS. X-ray diffraction (XRD) analysis indicates that no austenite phase was present in solutionized and quenched condition, but at overaged condition there may have 6 to 9% of volume fraction of reverted austenite [3]. It is reported that, copper precipitates produced during isothermal ageing are predominantly located at dislocations, where as during continuous cooling – dominant mode of precipitation was interphase precipitation, not restricted to be located at dislocations only, rather than finely spaced lines of precipitates that are effective at precipitation strengthening [11]. In context to it, it can be assumed, when diffusion bonding joining will be carried out at higher temperature and slowly cooled, no significant changes in strength has come for PHSS. The parent metals were machined to make samples of 30mm length and 15mm diameter. Conventional metallographic techniques were used to prepare mating surfaces of samples with 1µm diamond paste, subsequently cleaned in acetone and dried in air. The polished surface of PHSS and Cp-Ti were kept in contact in a jig. A load of 3.5 MPa was applied uniaxially along the longitudinal direction of the assemblies. The diffusion bonding was carried out at 800°C, 850°C, 900°C, 950°C and 1000°C for 3.6 ks in (4 to 6)×10¯3 Pa vacuum. During processing heating rate was 0.24°Cs-1 and the samples were allowed to cool in vacuum after processing at the rate of 0.1°C s-1. A transverse section was taken from the bonded samples and prepared by usual grinding and polishing techniques. The Cp-Ti and PHSS sides were etched with Kroll’s reagent ( 6 ml HCl, 2 ml HF, and 92 ml water ) and Fay’s reagent ( 40 ml HCl, 25 ml ethyl alcohol, 30 ml H2O and 5 gm CuCl2 ), respectively. The micro structure in the reaction zone was revealed in a light microscope (Correct ISDME TR5). The polished samples were also examined in a Scanning Electron Microscope ( LEO-S 440 ) using backscattered
(BSE) mode to exhibit finer structural details in the diffusion zones of the couples. Composition of the chemical species in the reaction layers as well as diffusion profile (Fig.3) cross their width was determined in atomic percent by Electron Probe Micro Analyzer (CAMACA Sx100) using 15 kV accelerating voltage, beam current 15nA and beam size 1 micron diameter. Phase width was also determined from SEM-BSE observation. Five such readings were taken from different location to find out arithmetic mean thickness of the layers. A set of bonded assemblies was fractured under shearing force and the phases on both sides of the fractured surfaces were identified in X-ray diffractometer ( Philips PW1710 ). A Co target was used at an operating voltage 40 kV, with sample current of 30mA and step height of 0.02° (=2θ) during diffraction study (800°C-950°C). The tensile testing of the diffusion bonds was carried out Static Instron Testing Machine of capacity 50KN ( Model 342) at a cross head speed of 0.00083 mm s-1 using sub-size cylindrical specimen as per ASTM (Vol.03.01 E8M-97). The original interface was at the centre of the gauge length of the tensile sample. The gauge diameter and the length of the sample were 4±0.05 and 20 mm, respectively. The shear strength of the bonded joints was evaluated at room temperature using a screw tensile testing machine set a crosshead speed 8.3 × 10-3 mm s-1. The shear test specimen were machined to a diameter of 10mm. Reproducibility of data was verified by repeating the evolution procedure for both tensile and shear testing. Micro-Hardness of interface layers were taken by micro-hardness testing machine (LICA-VMHT) and a micro-hardness profile (Fig.5) were obtained across the interface of different bands.
III. RESULT AND DISCUSSION
The optical micrograph (OM) of the diffusion bonds are shown in Fig:1. The interface of the transition joints is clearly visible without any discontinuity. The area ‘I’ indicates martensite phase with small volume fraction of ferrite of the parent PHSS. The zone contains heavily etched band ‘II’, which is the combination of different intermetallic layers, unetched stabilized β-Ti ( zone ‘III’ ) and α–β titanium ( area ‘IV’ ). Diffusion of strong β stabilizers element like Fe, Cr, Ni, and Cu from PHSS to titanium side retained high temperature phases of Ti even at room temperature in zone ‘III’. The area ‘IV’ indicates widmanstätten α–β titanium, containing the same β-stabilizers but in lesser quantity with respect to earlier one. These elements lower α–β eutectoid transformation
Table 1. Chemical compositions of the parent materials (wt%)
Alloy C Fe Ti Mn Si S P Cr Ni Cu Nb+Ta O N H cp Ti 0.02 0.10 Bal - - - - - - - - 0.15 0.02 0.0011 PH ss 0.044 Bal - 0.52 0.33 0.011 0.03 16.4 4.13 3.12 0.31 - - -
Table 2. Mechanical properties of base metals at room temperature
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temperature of Ti. During cooling β-Ti transforms to α–β titanium, acicular in nature, where the shaded α-Ti occured in the bright matrix of β-Ti [6,10].
(d)
(e)
Figure 1: Optical micrographs of the diffusion bonded joints processed for 3.6 ks at (a) 800oC (b) 850oC (c) 900oC (d) 950oC (e) 1000oC. The SEM-BSE images of the transition joints are shown in Fig.2. The joints formed at 800oC and 850oC clearly exhibit one distinct layer. This area is enriched with Ti(~51.2 – 55.4 at%) and Fe(~36.6 – 41.4 at%) with a small quantity of Mn (~0.3 – 0.4 at%) , Ni (~0.8 – 1.4 at%) ), Cu(~0.3 – 0.9 at%) and Cr ( balance); hence the compositions indicate the phase mixture of λ+FeTi phases. At 900oC joining temperature; in addition to λ+ FeTi phases layer, two new reaction layers have been observed at PHSS side with composition of Ti (~2.7 – 3.9 at%), Cr(~24.3 – 26.7 at%), Mn(~0.3 – 0.4 at%), Ni (~0.9 – 1.5 at%), Cu(~1.0 – 1.3 at%) and Fe (balance). Presumably, the area is phase combination of λ+α-Fe. Adjacent to λ+α-Fe phase mixture the bright area is enriched with Ti(~9.48 – 11.2 at%), Cr(~21.6 – 22.6 at%), Mn(~0.2 – 0.5 at%), Cu(~0.9 – 1.2 at%), and Fe(balance); hence the compositions indicate the phase mixture of λ+χ. In addition to these phases, the wide shaded region at the interface which is observed in between FeTi + β-Ti and λ+ χ phase-layers. The area is enriched with Cr (~5.5 – 6.5 at%), Ti (~27.1– 32.9 at%), Mn(~0.1 – 0.4 at%), Cu(~0.8 – 1.0 at%), depleted in Ni(~1.5 at%) and Fe (balance). This area is perhaps λ+FeTi phase mixture. It has been found that chromium enrichment occurs in the PHSS side. The α-Fe contains 35 at% of chromium in solid solution and migration of titanium from Cp-Ti side decreases the activity of Cr. Diffusion in that case occurs down to the
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activity gradient rather than concentration gradient to form a Cr rich layer[2,5,6,10]. The volume fraction of layers is not enough to resolve distinctly by SEM. So in Fig.2 no additional image effect has been observed. With the increase of temperature at 950oC; next to λ+FeTi phase mixture the shaded reaction band contains Ti (~52 – 53.9 at%) and Fe (~39.2 – 40.8 at%) in association with Cr (~4.3 – 4.4 at%), Ni (~1.4 – 1.6 at%) and Cu ( balance ), which is presumably the phase mixture of FeTi + β-Ti . This FeTi+β-Ti phase mixture has been observed and absence of this phase at below this temperature perhaps due to presence of low volume fraction.
(a)
(b)
(c)
(d)
(e)
Figure 2: SEM-BSE images of the diffusion couples processed for 3.6 ks at (a) 800oC (b) 850oC (c) 900oC (d) 950oC (e) 1000oC.
At subsequent increase of temperature at 1000°C, no further addition of reaction band has been indicated; rather the existing width of all reaction bands increased significantly which indicates the presence of increased volume fraction of brittle intermetallic phases. The width of the different reaction layers are furnished in Table 3. The extent of mass transfer depends on the processing temperatures. Increasing of joining temperature drives more number of atoms which migrate across the interface, that are responsible for the widening of the reaction layers. Assuming the growth of individual layer to be parabolic, as it is considered to be diffusion controlled and taking the overall thickness in consideration, the growth can be expressed by the following relations:
X2 = kt (1)
and 0kk = ⎟⎠⎞
⎜⎝⎛−
RTQexp (2)
Here, X is the thickness of the reaction layer (m); t is time of the bonding (s), T is bonding temperature (K); k is growth velocity of the reacting layer (m2s-1); k0 is growth constant (m2s-1),
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Table 4. Activation energy for the reaction layers
Reaction Layers
Activation energy ‘Q’(kJmol-1)
k0 (m2s-1)
α-Fe+λ 209.9 2.5 x 104
λ+χ 80.75 4.0 x 10-9
λ+FeTi 34.50 2.7 x 10-14
β-Ti 122.1 3.2 x 10-2
0 5 10 15 20 25 30
0
20
40
60
80
100
Temperature: 800oCTime : 3.6 ksAt
omic
per
cent
Distance in micrometer
TI CR FE NI CU
(a)
-5 0 5 10 15 20 25 30 35 40
0
20
40
60
80
100
Temperature: 850oCTime : 3.6 ksA
tom
ic p
erce
nt
Distance in micrometer
Ti Cr Fe Ni Cu
(b)
-5 0 5 10 15 20 25 30 35 40
0
20
40
60
80
100
Temperature: 900oCTime : 3.6 ksA
tom
ic p
erce
nt
Distance in micrometer
TI CR FE NI CU
(c)
Q is activation energy for layer growth (kJmol-1); R is real gas constant (8.314 Jmol-1K-1) The plot of ln(X) vs 1/T and the linear fit of the same give the values of Q from the slope and the value of k0 from the intercept. Activation energy (Q) and growth constant (k0) values of α-Fe+λ, λ+χ, λ+FeTi and β-Ti phase or phase mixture are given in Table-4. The growth rate of α-Fe+λ phase is faster than that of the other phase or phase mixture.
Fig.4 shows the X-ray diffraction analysis, which confirmed the presence of different intermetallics in the fracture interface. The X-ray diffraction data indicated the formation of reaction products of α-Fe, σ, χ, Fe2Ti, FeTi, Cr2Ti and β-Ti. The λ phase is the solid solution of Fe2Ti and Cr2Ti, and is identified separately. The reaction zone does not show existence of any Cu or Ni bearing intermetallics. The presence of σ phase is confirmed by X-ray diffraction analysis. The absence of σ phase in SEM-BSE micrographs may be due to its small volume fraction and /or finer size. σ phase contains principally Fe and Cr, and is known to be a brittle phase. Ti migration in the stainless steel decreases the activity
Table 3. Width of the intermetallic phases formed in the reaction zone
Bonding temperature (°C)
Width of the reaction products in the diffusion zone (µm)
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of Cr. Hence, diffusion occurs down the activity gradient: rather than the concentration gradient. Cr enrichment occurs by uphill diffusion of the same and during cooling the chromium enriched region transforms to σ phase [2,5,6,10]. The variations of the mechanical properties of bonded joints with the change in bonding temperature are given in Table.5 and Fig.6. It can be seen that bonding temperature at 800°C and 850°C, both tensile and shear strength are minimum due to lack of contact between the mating surfaces, as the yield stresses of the materials still remain high. Low processing temperature also leads to a minimum thermal excitation and the extent of diffusion of alloying elements is limited at both interfaces. Moreover, diffusion of alloying elements promote intermetallics formation in the diffusion zone which lower the bond strength. With increasing the bonding temperature at 900°C, tensile and shear strength as well as ductility of the bonded sample is increased due to better extent of promotion of same phenomena, lack of which the bond strength is found low at bonding temperature 800°C to 850°C. Maximum strength is achieved at 950°C due to increasing contact area between the mating surfaces subsequently the plastic collapse of the faying surface asperities that augment and also promote better inter-diffusion and which, in turn gives better alloying at diffusion zone in addition to it, increment of the width of the reaction bands at the diffusion zone which indicate the increase of intermetallics volume fraction. However, these volume fractions are negligible compare to temperature 900°C; so the brittle intermetallics affects to the bond strength not much accountable here. It has been also found, that at 1000°C, bond strength dropped owing to increased volume fraction of brittle intermetallics. This phenomenon overbalances the gain in strength due to increased atomic thermal excitation. From Table 3, it is evident, that total width of intermetallic compounds at 1000°C joining temperature is five times larger than the total width of the intermetallic compounds-layer at 900°C joining temperature and two and half times larger for the sample processed at 950°C. In correlation with the micro-hardness value (Fig.5), it has been found, higher the bonding temperature higher is the micro-hardness value which may apparently oppose the previous experimental results of lower bond strength at 1000°C in spite of better inter-diffusion of alloying elements and subsequent alloying. However, indeed higher hardness value comes from generation of strain due to poor accommodating of excess volume fraction of brittle intermetallics at above
20 30 40 50 60 70 80
500
1000
1500
2000
2500
σ, C
r 2Ti
α-Fe
, χ
FeTi
, χC
r 2Ti
α-Fe
, Fe 2Ti
σ, F
e 2TiCr
2Ti
χ
800 oC
850 oC
950 oC
900 oC
Inte
nsity
(arb
itary
)
o2θ
(a)
20 30 40 50 60 70 80200
400
600
800
1000
1200
1400
1600
α-Ti
α-Tiχ
α -Ti
, β-T
i
Cr2Tiα-
Ti
β-Ti
α-TiFe
TiCr
2Ti, F
e 2Ti
χ Cr2Ti
Fe2Ti
Cr2Ti
, FeT
iα-
Ti
α-Ti
800 0C
900 0C
950 0C
850 0C
Inte
nsity
(arb
itary
uni
t)
o2θ
(b) Figure 4: X-ray diffraction study for the diffusion bonded joints processed at (a) 800oC (b) 850oC (c) 900oC (d) 950oC for 3.6 ks, (a) PHSS side, (b) Ti side.
-20 -15 -10 -5 0 5 10 15 20 25
150
200
250
300
350
400
450
500
550
600
650
700
750
Mic
roha
rdne
ss (H
V)
Distance in micrometer
D800oC D850oC D900oC D950oC D1000oC
Time: 3.6 ks
Figure 5: Microhardness profile at and near bonding interface processed for 3.6 ks at (a) 800oC (b) 850oC (c) 900oC (d) 950oC (e) 1000oC
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800 850 900 950 1000100
150
200
250
300
350
Shear Strength Tensile Strength Breaking Strain
Temperature (oC)
Stre
ngth
(MPa
)
2
4
6
8
10
12
14
Stra
in (%
)
Figure 6: Mechanical Properties of the diffusion bonded joints processed at (a) 800oC (b) 850oC (c) 900oC (d) 950oC (e) 1000oC for 3.6 ks. 950°C, the same finally gives the lower bond strength with poor ductility.
Now, comparing the theory for bonding temperature 950°C; here due to well optimisation of accommodating of brittle intermetallics with much lower volume fraction of the same, no such noticeable strain generation has been obtained rather it shows enhance ductility. More over, at 900°C and 950°C bonding temperature; the obtained hardness value is same though the higher bonding strength is obtained at latter temperature. In a close look at Table 3 no such significant changes in reaction products and its width have been observed except apparently absence of FeTi + β-Ti but it suppose to be due to low volume fraction unable to detect by SEM-BSE. The whole reason indicates to insufficient inter-diffusion of alloying elements and relatively lesser extent of collapsibility of mating surfaces at lower processing temperature at 900°C. As the temperature increases to 950°C the above lacks just well compensated togetherly and enhance the bond strength. VI. CONCLUSIONS 1. The optimum parameter for diffusion bonding of commercially pure titanium and precipitation hardening stainless steel was found to be : temperature 950°C, time 3.6 ks, under the 3.5 Mpa uniaxial load in (4 to 6)×10–3 Pa vacuum. The tensile strength and sear strength of the joint achieved (344.3 ± 4) Mpa and(260.1± 4) Mpa respectively as well as 12.8% breaking strain. 2. When commercially pure titanium and precipitation hardening stainless steel were diffusion bonded at optimum parameter, the layer of intermetallic phases like λ+α-Fe (~15.1µm), λ+χ (~1.2µm), λ+FeTi (~1.4µm), FeTi+β-Ti (~ 1.1µm) and β-Ti (~61.1µm) are formed at the transition zone of the joint interface and obtained highest tensile and shear strength with a good breaking strain. 3. Lower and higher bonding temperatures than optimum temperature gave weak joints, due to lack of interdifussion of each base metal and excess interdiffusion of the same with increased intermetallic phases as well as generation of internal stress respectively.
ACKNOWLEDEMENT The author would like to thank to Dr. S.Chatterjee, Professor, Department of Metallurgy and Materials Engineering, Bengal Engineering and Science University Shibpur, Howrah, India, for his kind supervision.
REFERENCES
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[7] He P, Jhang J, Zhou R, Li X. Diffusion BondingTechnology of a Titanium Alloy to a Stainless Steel Web with an Ni interlayer. Material Characterization 1999;43:287–292,
[8] Kato H, Abe S, Tomizawa T. Interfacial Structure and Mechanical Properties of Steel-Ni and Steel –Ti Diffusion Bonds. J Matter. Sci. 1997;32:5225 – 32.
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[10] M. Ghosh, A. Laika, K. Bhanumurthy, G. B. Kale, J. Krishnan and S. Chatterjee, Evolution of Interface Microstructure and Strength Properties in Titanium – Stainless Steel Diffusion Bonded Transition Joints. 1578 Mat Sci. and Technol.. Dec’2004, Vol.20.
[11] R. H. Wagoner, Smith Chair ( revised by) Physical Metallurgy of Steel, MSE 661.
About the author: The author completed his Master of Engineering in Metallurgy and Materials Engineering from Bengal Engineering and Science University, Shibpur, Howrah, India, in the year 2007. He has 12 years of industrial experience at Rifle Factory Ishapore, India. The next mission is to enter into a research program that leads to Ph.D. degree.
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