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Diffusion bonding of Q345 steel to zirconium using an aluminum interlayer
Journal of Materials Processing Tech. 275 (2020) 116352
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Diffusion bonding of Q345 steel to zirconium using an aluminum interlayer a
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Yang Zhang , Bangzuan Long , Kai Meng , Aleksandr Gohkman , Ye Cui , Zhongwu Zhang
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Key Laboratory of Superlight Materials and Surface Technology, Ministry of Education, College of Materials Science and Chemical Engineering, Harbin Engineering University, Harbin, 150001, China Department of Physics, South Ukrainian National Pedagogical University, 65020, Odessa, Ukraine
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ARTICLE INFO
ABSTRACT
Associate Editors: C.H. Caceres and S.J. Na
The microstructure, interface reactions and mechanical properties of the Q345 steel/Al/Zr bonding interface were investigated. The thicknesses of the reaction layer in both Al/Fe and Al/Zr interfaces are temperaturedependent, with a higher temperature yielding a larger thickness. Both experimental results and thermodynamic calculation confirmed that the diffusion transition regions near the Al/Fe and Al/Zr interfaces mainly consist of Al5Fe2 and Al3Zr with thin Al3Fe and Al3Zr2 transitional phases, respectively. The growth kinetic equations of Al5Fe2 and Al3Zr as a function of temperature were investigated. The fitted curves from the equations well agree with the experimental results. The hardness of Al5Fe2 is higher than Al3Zr and the base materials. The propagation of shear crack occurs within Al5Fe2. Optimized shear strength can be obtained by controlling the thickness of reaction layer. The grain growth and fracture mechanisms of Al5Fe2 and Al3Zr at the interfaces were also discussed based on electron back scattering diffraction (EBSD) analysis of both reaction layers and fracture morphology.
Keywords: Diffusion bonding Microstructural evolution Growth mechanism Mechanical properties Reaction layer
1. Introduction Srikanth et al. (2017) reported that zirconium alloys are widely applied as the in-core structural material in various reactor vessels due to the superior properties. Whereas, stainless steels are Srikanth et al.usually used as the out-of-core components because of the good corrosion resistance. Therefore, bonding between steel and Zr-based alloys are often performed at various thermal reactors. However, Shaaban et al. (1978) found that direct joining of Zr-based alloys and stainless steel has encountered severe problems due to the different physical and chemical properties, such as the melting temperature and thermal expansion coefficient, which generate residual stresses leading to the microcracks formation at the bonding interface. Oliveira et al. (2016) indicated recently that the formation of brittle intermetallics (ZrFe2, ZrFe3 and ZrFe4) together with the crystallographic mismatch between Zr and Fe can reduce the bonding strength and lead to brittle failure. Oliveira et al. (2019) used friction element welding to keep the maximum welding temperature below the liquidus temperature of the high strength steel and friction element, which lower the risk in the formation of brittle intermetallic compounds. Ahmad et al. (2003) used electron beam welding to join Zircaloy-4 and stainless steel, although defects like porosity, voids and cracks were avoided, the harmful intermetallic compound Zr(Cr,Fe)2 formed. Zhang et al. (2019)
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investigated the interfacial microstructure between Mg and 22MnB5 boron steel using resistance spot welding, and found that the source of crack was due to the formation of brittle Mg-Al phase. Bhanumurthy et al. (2001) found that the diffusion reaction was extremely sluggish between 403 stainless steel and Zr-2.5 wt% Nb alloy, and diffusion bonding could be conducted only at a high pressure and temperature (10 Mpa, 900 °C). Therefore, studies on the bonding of steel and zirconium should be performed to minimize or prevent the brittle phases formation to obtain a good bonding. A simple and feasible method for this aim is using an interlayer. Akhter at al. (2005) proposed that the interlayer can prevent brittle phases formation at the interface, and is helpful for reducing the bonding temperature and pressure. Previous studies used Cu, Ni, Ti, and Ta as interlayers to avoid the generation of brittle intermetallics between Fe and Zr. For example, Aboudi et al. (2017) found that Cu interlayer failed to restrict the formation of brittle phases but was benefit to reduce the hardness of the reaction region. Meng et al. (2018) studied the thickness effect of Cu interlayer on the microstructure evolution, shearing and bending properties. The optimized shearing and bending strength can be obtained by controlling a suitable Cu interlayer thickness. Srikanth et al. (2017) used Ni and Ti interlayers to join 304 L stainless steel to Zircaloy-4 alloy, and found that when temperature reached 850 °C and 900 °C, the joints would fail owing to the cracks
Corresponding author. E-mail address: [email protected] (Z. Zhang).
https://doi.org/10.1016/j.jmatprotec.2019.116352 Received 11 April 2019; Received in revised form 25 July 2019; Accepted 28 July 2019 Available online 30 July 2019 0924-0136/ © 2019 Elsevier B.V. All rights reserved.
Journal of Materials Processing Tech. 275 (2020) 116352
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formation and propagation through the reaction layer (RL). Various methods have been conducted to join steel and zirconium, such as solid-solid diffusion bonding, explosion welding, fusion welding, and electron beam welded joints. These methods like explosion welding, however, have drawbacks of either high cost or lacking versatility. Perona et al. (1966) revealed that brittle intermetallic compounds, such as ZrFe2, ZrFe3 and Zr(Cr,Fe)2, were formed in the welded zone when using fusion welding joined stainless steel and Zircaloy-4These compounds are harmful to the mechanical property and corrosion resistance of the welded joints. However, many disadvantages of welded joints can be overcome using vacuum diffusion bonding process. Compared to the welding process, Springer et al. (2011) revealed that the base materials were not melted in diffusion bonding, which could avoid the residual stress of interface region in the joints caused by thermocycling. In the present study, the aluminum interlayer is used to bond Q345 steel and pure zirconium using diffusion bonding process. Kang and Kim (2015) have conducted on the joining of Al/Fe and Al/Zr. Their results confirmed that the diffusion bonding can provide a good strength. The choice of the base materials in this study are mainly due to the following reasons:
et al. (2007) found that the surface roughness had a large effect on the bonding, and thus all the base materials were polished carefully before bonding. Sand papers with various roughness of 800, 100, 1500, 2000, and 3000 were successively used to polish to a final roughness of about 1 μm. The polished base materials were cleaned by water first and further by absolute alcohol for 15 min in Ultrasonic instruments, followed by blowing dry with air blower. Diffusion bonding was performed from 530 °C to 620 °C with a pressure of 3 Mpa for 3 h in a vacuum hot press sintering furnace. The vacuum of the furnace is less than 1 × 10−3 MPa. A more detailed sample information and experimental procedure can be found in our previous study (Meng et al., 2018). The bonded samples at different temperatures were cut to sheets with a section of 12 × 12 mm2, and then the bonding surfaces were polished using mechanical polishing for microstructure observation and hardness measurement. At least five indentations (along the interface direction) were measured for each distance to obtain an average value. The microstructures were observed by metallographic microscope (LEICA DMIRM, Germany) and an FIE-Quanta 200FEG scanning electron microscope (SEM) with back scattered mode (BSE). The acceleration voltage is 20 kV. The phase compositions of the joint interfaces were identified by line scanning of energy dispersive spectrometer (EDS, FIE, Quanta 200FEG). X-ray diffraction (XRD) was performed at Rigaku D/max-TTR- III. To identify the products by XRD in the interface, the bonded samples were stripped layer by layer along the bonding interface to make sure the reaction products were exposed thoroughly. The reaction products at Zr and Fe sides were determined by XRD, separately. The microhardness of the bonding interface were conducted by an HVS-1000Z tester with 10 g load for 25 s. Shear test was performed using a San-Si tester. The schematic diagram of shear test can be found in our previous paper (Meng et al., 2018). A compression speed of 1 MPa/s was used. According to Yan et al. (2010), the shear stress (τ) was calculated from the following relations:
1 Q345 steel belongs to low-carbon high-strength steel which also has good plasticity and welding performance. The content of alloy elements is relatively low, which can avoid the effects of other alloy elements on the diffusion reaction between Fe and Al. 2 Pure zirconium has very low impurity content, and is easy for the mechanism investigation of diffusion bonding between Zr and Al. 3 Aluminum is chosen as an interlayer due to its relatively low melting point (660 °C) over other metals and, therefore, critical conditions for diffusion bonding can easily be met. Although many studies have been conducted on the joining of Al/ Fe, the interdiffusion and growth mechanism of interface remain contradictory. The most common view is that a large number of vacancies facilitate the fast diffusion of Al atoms, leading to the growth of Al2Fe5 along [001] direction (c axis) (Takata et al., 2015). However, the joining between Al and Zr using diffusion bonding is rarely investigated. Only interdiffusion and reaction were reported but the growth mechanism of interface remained unknown. Therefore, the diffusion bonding between steel and zirconium using an aluminum interlayer is of great importance for further understanding the interface growth mechanism. Mehrer (2007) found that diffusion in solids strongly depends on temperature, being slow at low temperatures but appreciable at high temperatures. Pressure also plays an important role, but is far less striking than temperature. Travessa et al. (2002) revealed that chemical reactivity at the interface is also a significant requirement for bonding, as it enables the occurrence of adhesion at reasonably low temperatures and moderate external pressures. The experimental conditions in this study can be identified as low temperature and low pressure incomparison to the common ones. The bonding at high temperature is easy to cause hydrogen absorption and lead to hydrogen embrittlement. Srikanth et al. (2017) pointed out that the thermal mismatch and intermixing can be limited by decreasing the bonding temperature. In this paper, effects of temperature on the microstructure, microchemistry of the interface, kinetic and thermodynamic features, hardness and shear properties are investigated in detail. The fracture mechanism, growth mechanism of the reaction phases at the interfaces are also discussed based on the fracture morphology and EBSD analysis of both reaction layers.
=
P ae
(1)
Where a and e are the width and height of test interface, respectively. In this study, a and e are 25 mm and 3 mm. The growth orientation law of Al-Fe and Al-Zr interfaces were studied by electron backscatter diffraction, EBSD (Hitachi S3400 scanning electron microscope equipped with HKL Channel 5 software). Fig. 1 shows the schematic diagram of coordinate system (RD-TD-ND) in EBSD measurements. 3. Results and discussion 3.1. Microstructure evolution Fig. 2 shows the interfacial microstructure evolution between 530 °C and 620 °C. The thickness of Fe/Al diffusion layer is almost zero at
2. Experimental procedure Pieces of Q345 steel (100 × 100 × 6 mm3) and pure zirconium (100 × 100 × 1.5 mm3) were used for vacuum diffusion bonding. A 500 μm thick pure aluminum foil was selected as an interlayer. Chen
Fig. 1. Schematic diagram of EBSD test coordinate system. 2
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Fig. 2. Microstructure of the bonding interface at (a) 530 ℃, (b) 545 ℃, (c) 560 ℃, (d) 575 ℃, (e) 590 ℃, and (f) 620 ℃.
530 °C, indicating that almost no reaction occur at this temperature. The thicknesses of both Fe/Al and Zr/Al reaction layers (RLs) increase as temperature increases, but with different changes. The increase in bonding temperature facilitates the atoms to diffuse across the interface, yielding a thicker reaction layer. The RL width of Al-Fe is almost 3 times wider than that of Al-Zr at 560 °C, indicating a higher reaction speed of Al-Fe than Al-Zr. From Fig. 2(a) to (f), one can observe that both interfaces develop towards interlayer Al owing to a larger diffusion coefficient of Fe and Zr to Al than that of Al to Fe and Zr at the same temperature (Fisher et al., 2013). When the bonding temperature reaches 590 °C, micro-crack emerges in the Al-Fe RL region. According to the binary phase diagram of Al-Fe, the Al-Fe compounds are all hard and brittle phases. In addition, Wang et al. (2014) reported that the thermal expansion coefficient of Al (23.5 × 10−6 K-1), Fe (12.2 × 10−6 K-1), and Al-Fe intermetallic compounds (18.94 × 10−6 K-1 for Al5Fe2, 19.68 × 10−6 K-1 for Al3Fe) are quite different. Thus, the microcracks are observed within Al-Fe intermetallic compound layer. With the further increase of temperature to 620 °C, almost all the Al interlayers have reacted with Fe and Zr, and large cracks emerge between the products of Al-Fe and Al-Zr. The thickness of Al-Zr products is larger than that of Al-Fe, indicating that bonding temperature has a totally different effect on the growth rate. Kidson and Miller (1964) also observed many voids due to the diffusion imbalance across the interface due to Kirkendall effect. The voids are near the Zr
side due to the higher diffusion coefficient of Zr to Al than Al to Zr. 3.2. Phase identification To analyze the elemental compositions across the reaction interface at 545 °C, 560 °C and 575 °C, EDS line scanning was performed and shown in Fig. 3. The element composition distributions are shown in Table 1. According to the atom ratio of Al/Fe and Al/Zr on the surfaces, the reaction products (Al5Fe2 and Al3Zr) were identified. Al5Fe2 exhibit a tongue-like morphology with highly irregular peaks and valleys orientated towards the iron. In the reaction of Fe and Al, Li et al. (2016) found that Al5Fe2 is always easily formed whether in solid phase reaction or solid-liquid reaction, and exhibit a morphology of wave shape towards the Fe direction. Dickson et al. (2014) found that thick Al3Zr reaction layers formed in the diffusion bonding of Al and Zr in the temperature range from 425 to 625 °C, and dense tiny voids were observed on the side close to Zr due to Kirkendall effect. Although this effect also exist in 545–575 °C, it was not observed because the atom diffusion was relatively slow in this temperature range. Only when the temperature reaches 620 °C, large voids form within Al3Zr due to the aggregation of vacancies as shown in Fig. 2f. In many binary diffusion couples, in addition to the formation of a stable phase, there will be one or more transition phases between the stable phase and the diffusion couple. The transition phases formed at 3
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Fig. 3. The chemical compositions across the bonding interface of steel/Al/Zr at (a) 545 ℃, (b) 560 ℃, (c) 575 ℃.
3.3. Phase formation sequence in Fe/Al and Zr/Al binary solid phase diffusion systems
Table 1 Elemental compositions of Al-Fe and Al-Zr reaction productions. T(oC)
545 560 575
Q345 steel side
Zr side
Fe(at.%)
Al(at.%)
Al/Fe
Phase
Al(at.%)
Zr(at.%)
Al/Zr
Phase
27.0 27.3 27.8
73.0 72.7 72.2
2.7 2.7 2.6
Al5Fe2
73.9 73.7 74.0
26.1 26.3 26.0
2.8 2.8 2.8
Al3Zr
The driving force of phase formation is mainly determined by the change of Gibbs free energy ΔG0, which can be expressed as
G0 = H 0
(2)
T S0
Where T is Kelvin temperature, and are the change in enthalpy (heat) and entropy, respectively. In solid-state reaction, Lee et al. (2003) revealed that the entropy change is further smaller than the enthalpy change, and thus it can be negligible. Then G 0 can be approximated by H 0 , indicating that the phase formation can be predicted by the heat of formation. Pretorius (1990) proposed that the formation law of products during solid phase diffusion can be well predicted according to EHF (effective heat of formation) theory. Pretorius et al. (1993) state that many factors such as impurities, lowest eutectic, atomic mobility, diffusing species, could affect the actual concentrations that are available for interaction at the growth interface. Therefore, the term effective concentration is used in the effective heat of formation concept. The released heat is dictated by the effective concentration of the limiting element (Pretorius et al., 1991). Therefore, the effective heat of formation can be described as follows:
H0
the junction of Al5Fe2, Al3Zr and the parent metal or interlayer. A high magnification SEM image of the RLs at 560 °C is shown in Fig. 4. The micron transition phases can be observed at the surfaces of Al5Fe2/Al and Al3Zr/Zr. The major elemental compositions at each spot within the transition phases detected by EDS show that the transition phases can be identified as Al3Fe and Al3Zr2, which has also been observed by Chen et al. (2016). The average thickness of transition phases also increase with the increase in temperature. The layered (Al3Fe) and tongue-like (Al3Zr2) morphology become more and more clear with the increase in temperature. The thickness of Al3Fe layer is much thinner than Al5Fe2 layer, which can be attributed to the much lower diffusion coefficient (DAl5Fe2 = 1.84 × 10−8 m2/s and DAl3Fe = 9.2 × 10-10 m2/s) (Bouché et al., 1998). The phase components of Fe/Al and Zr/Al interfaces were further verified by XRD (Fig. 5). From the patterns, it is found that the reaction product mainly consists of Al5Fe2 at steel side, and Al3Zr is observed on the zirconium side, which is well agreed with the EDS results and the following thermodynamic calculation results. However, the secondary transitional products, Al3Fe and Al3Zr2, are not detected due to the detection limit of XRD technology. The individual small unrecognized peaks maybe due to the oxide formation of Fe2O3 and Al2O3.
H = H0 ×
S0
effective concentration limiting element compound concentration limiting element
(3)
Theron et al. (1996) showed the formation enthalpy of the binary compounds in Al-Fe and Al-Zr systems. The lowest liquidus (eutectic) temperature is 660 °C according to the Al-Fe binary phase diagram, and the effective concentrations of Al and Fe at the interface are 99.1% and 0.9%, respectively. Thus, Fe is a limiting element. The effective heat of formation of different Al-Fe phases are shown in Fig. 6a, and it is obvious that ΔHʹAl5Fe2 < ΔHʹAl3Fe < ΔHʹAl2Fe < ΔHʹAlFe3. Al5Fe2 has the lowest effective enthalpy of formation, and Al3Fe takes second place, indicating Al5Fe2 and Al3Fe are the first and the second products, respectively, which further confirms the experimental results. According
Fig. 4. A high magnification SEM image of the RLs, (a) Al5Fe2/Al interface and (b) Al3Zr/Zr interface. 4
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Fig. 5. XRD patterns of (a) Fe/Al and (b) Zr/Al interfaces.
to EHF model, after the formation of the first product in binary metal system, the second product will form at the interface between the first product and the excess metal (M) and enrich the M element, and Pretorius (1990) confirmed this phenomenon in the experiments. Similarly, the lowest temperature on the liquidus curve is 660 °C in the Al-Zr binary system, and the effective concentrations of Al and Zr at the interface are 98.0% and 2.0%, respectively. In this case, Zr is considered as a limiting element. The effective enthalpy of formation of Al3Zr, Al2Zr, Al3Zr2 and AlZr were calculated and shown in Fig. 6(b). It is found that ΔHʹAl3Zr < ΔHʹAl2Zr < ΔHʹAl3Zr2 < ΔHʹAlZr, where Al3Zr has the lowest effective formation enthalpy of -3.28 kJ·(mol·at.)−1 being the first reaction product, which agrees with the experimental results. However, the calculated second product is Al2Zr rather than Al3Zr2 (experimental result). This phenomenon is named as “Phase jump behavior”, and has also been observed in Al-Pd (Colgan, 1987), Al-Pt (Hultgren et al., 1973) and Al-Au (Campisano et al., 1975) binary system. Laik et al. (2004) calculated the Gibbs free energy of Al2Zr or Al3Zr2 on the interface of α-Zr(Al)/Al3Zr by the solid solution model of binary system. The results show that ΔGAl3Zr2=-10.77 kJ/mol and ΔGAl2Zr=-7.18 kJ/mol, indicating that Al3Zr2 has a larger driving force and is preferentially formed between α-Zr(Al)/Al3Zr.
Table 2 The average thicknesses of the RLs. Temperature (oC)
Al5Fe2 (μm)
Al3Zr (μm)
545 560 575 590
84 93 154 191
20 31 65 124
diffusion temperature. Based on the experimental data at temperatures of 545 °C, 560 °C and 575 °C in Table 2, the relationships between lnx and T for Al5Fe2 and Al3Zr can be fitted as:
lnx = 9.22
15.25 ×
lnx = 25.78
30.00
1000 (Al55Fe2) T
(5a) (5b)
Then, the growth kinetic equations of Al5Fe2 and Al3Zr can be described as
xAl55Fe2 = 1.01 × 10 4 exp
126800 , QAl55Fe2 = 126.8kJ•mol RT
249400 , QAl3Zr = 249.4kJ •mol RT
1
(6a)
3.4. Growth dynamics of Al5Fe2 and Al3Zr
xAl3Zr = 1.57 × 1011 exp
The measured average thickness of the RLs from SEM-BSE are shown in Table 2. Luo and Acoff (2004) discovered the thickness of reaction layer, x, can be expressed as
The fitted and experimental temperature dependences of the Al5Fe2 and Al3Zr layer thickness are plotted in Fig. 7. The thicknesses of RLs from the calculations are well identified with the experimental results (the experimental result of Al5Fe2 thickness is zero and thus not shown here), which can be observed from the inset of Fig. 7. With the further increase in temperature, the growth curve of Al5Fe2 and Al3Zr intersects, and then the growth rates of these two phases begin to be reversal. From Fig. 2(e) and (f), one can observe that the thickness of Al5Fe2 is larger than Al3Zr layer at 590 °C, however, the thickness
lnx = ln (K 0 t )
Q RT
(4) −1/2
where K0 is a temperature dependent reaction constant (m·s ), t is diffusion time (s), R is Boltzmann constant (8.314 J·mol-1 K-1), and T is
Fig. 6. The effective heat of formation diagram, (a) Fe-Al and (b) Zr-Al. 5
1
(6b)
Journal of Materials Processing Tech. 275 (2020) 116352
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interface reaction process and the growth rate is controlled by the diffusion of the slowest element. As the diffusion activation energy of Fe atoms in aluminum was less than that of Al atoms in the steel like mentioned above, and the Al-Fe interface grows towards to Al interlayer shown in Fig. 2, thus the growth of the Al5Fe2 phase can be assumed to be controlled by diffusion of Al. Abbasi et al. (2016) observed the presence of Al3Fe phases along the grain boundaries of Al, which offer fast diffusion paths for Fe atoms to reach low-Fe areas and nucleate. However, Al3Fe phase was not observed from the EBSD results due to the thin thickness. Fig. 9 shows the grain orientation distribution across the interface of RD, TD, and ND directions. Here, we mainly focus on the orientation distribution of Al5Fe2, for which red represents [001] orientation, blue represents [100] orientation and green represents [010] orientation. It is obvious that the [001] orientation of most Al5Fe2 grains is along TD direction from Fig. 9(b). The grains parallel to RD and ND directions have two types from Fig. 9(a) and (c): the blue ones with [100] orientation and the green ones with [010] orientation. To further investigate the orientation distribution of Al5Fe2 grains, the full and inverse pole figures are measured (Fig.10a and 10b). Ding et al. (2018) and Takata et al. (2015) used the similar method to illustrate the orientation of η-Fe2Al5 phase. It can be seen from Fig.10 that [001] concentrated along TD direction. The elongated grain growth of Al5Fe2 has a strong orientation, that is, the c-axis [001] direction is perpendicular to the interface between steel and aluminum, which is also confirmed by Murakami et al. (2004). This kind of grow priority guarantees that the Al5Fe2 phase grows much faster than other intermetallic phases. The grain boundaries distribution of Fe-Al5Fe2-Al system are shown in Fig. 11(a) and (b). In general, Onuki et al. (2013) and Zhang et al. (2018) defined that the grain boundary angle of 2°-15° belongs to a small-angle grain boundary (green line) and a grain boundary greater than 15° is defined as a large-angle grain boundary (black line). The focus is the distribution of grain boundaries within Al5Fe2, in which the large-angle grain boundary accounts for 70%, mainly distributed near the Al side; the small-angle grain boundary accounts for 30%, mainly distributed near the Fe side. Small angle grain boundary density is the highest in the area contacting with Fe, which suggests that the strain due to the transformation from α-Fe phase to Al5Fe2 introduces a high density of dislocations into the neighboring α-Fe phase, developing a dislocation substructure around the interface (Takata et al., 2015). The large Al grains mainly consist of small-angle grain boundaries, while the grains within Fe are mainly large-angle grain boundaries. Therefore, there is a grain boundary angle difference at the junction of Fe to Al5Fe2 transition, and the small angle grain boundary position has a higher dislocation density. Burkhardt et al. (1994) reported Al5Fe2 is an orthorhombic crystal structure with a space group of Cmcm (a = 0.766 nm, b = 0.642 nm, and c = 0.422 nm). While α-Fe belongs to a body-centered cubic crystal structure with space group IM3M (a = 0.293 nm). Since the number of atoms in Al5Fe2 and α-Fe unit cells
Fig. 7. The fitted and experimental curves between the thickness of RLs and temperature. The inset shows a zoom from 790 K to 870 K.
reverses at 620 °C. The interface microstructure at 620 °C can be used to explain this phenomenon shown in Fig.2(f), from which we can observed that the Al intermediate layer is almost consumed by Fe and Zr, and the thickness of Al3Zr is far greater than Al5Fe2. When the temperature reaches 617.7 °C, the reaction rate flips up, resulting in the thickness of Al3Zr begins to be larger than Al5Fe2 (Fig. 7). The activation energy for Al3Zr is 249.4 kJ mol 1, which is much higher than Al5Fe2 (126.8 kJ mol 1), indicating that the reaction of Al5Fe2 is easier than Al3Zr. This can be reflected by the RLs thicknesses of Al5Fe2 and Al3Zr as shown in Table 2. 3.5. Growth orientation law of Al5Fe2 phase In order to investigate the growth orientation law of interface phases during diffusion bonding, EBSD was used to characterize the microscopic crystallographic characteristics of Al5Fe2 and Al3Zr. The phase distribution of Al-Al5Fe2-Fe region are displayed in Fig. 8(a), where red presents Al, yellow presents Al5Fe2 and blue stands for Fe. One can see that Al and Fe have larger crystal grains than Al5Fe2. The grain size distribution of Al5Fe2 is shown in Fig. 8(b), and exhibits an obvious trend: the smaller grains are mainly concentrated at the Al side; the larger and elongated grains are more close to the Fe side, which agree with the previous study (Abbasi et al., 2016). It can be inferred that the crystal nucleation occurs at the Al side, and the fine crystal grains are mutually swallowed and gradually grow towards the Fe side with the increase of diffusion time. Laik et al. (2013) showed that only when the stoichiometric composition of phase layer is satisfied, the phase layer grows, and its growth process behaves as a series of
Fig. 8. (a) The phase distribution of Al-Al5Fe2-Fe region, (b) grain size distribution of Al5Fe2. 6
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Fig. 9. The grain orientation distribution of Al-Al5Fe2-Fe, (a) RD, (b) TD, (c) ND.
Fig. 10. The full (a) and inverse (b) pole figures of Al5Fe2 grains.
are different, leading to an apparent volume expansion originating from the transformation from the α-Fe to Al5Fe2 (Takata et al., 2015). Richards et al. (1994) points out that there is a large difference in thermal expansion coefficient, leading a difference in the stress at the transition interface of Fe to Al5Fe2. The Al5Fe2 grains suffer from stress during grain growth, causing that the grain boundaries perpendicular to the stress are stretched and those parallel to the stress are compressed. That is, the c-axis of the
grain growth direction is subjected to tensile stress, while a and b axis perpendicular to the grain growth direction are subjected to compressive stress. The grain of Al5Fe2 is elongated towards the c-axis, indicating that the stress in this direction is relatively small. Thus, the transformation from Fe to Al5Fe2 becomes easier. However, the stresses in a and b directions are large, leading to an obstruction of the growth direction in these directions. It is found that the crystalline defects of Al5Fe2 possess 30% vacancies along the c-axes ([001] direction), which 7
Journal of Materials Processing Tech. 275 (2020) 116352
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Fig. 11. (a) and (b) are grain boundaries distribution of Fe-Al5Fe2-Al system.
offers a rapid diffusion path for Al5Fe2 and increases its growth rate, causing the formation of columnar Al5Fe2 and the corresponding residual steel substrate (Cheng and Wang, 2009). The growth of Al5Fe2 is controlled by the vacancy mechanism, and the grain boundary itself is the position of the vacancy source and the annihilation. The formation energy of the grain boundary vacancy perpendicular to the stress direction is low, and the number of vacancies is large; while the grain boundary vacancy formation energy of parallel stress is high, the number of vacancies is less, thus causing a certain vacancy concentration difference inside the grains. Therefore, a great quantity of vacancies gather in the c-axis direction, providing a favorable channel for the diffusion of atoms. As a result, the Al5Fe2 grains appear a slender morphology and the overall interface exhibits as a “wavy” shape.
lines represent high-angle grain boundaries. Then we mainly focus on the grain boundaries distribution of Al3Zr, and the misorientation angle distribution of Al3Zr is shown in Fig. 15(b). The large-angle grain boundaries possess 92.8%, and only a small part of small-angle grain boundaries are distributed in the grown up equiaxed grains. 3.7. Mechanical properties The mechanical properties of the steel/Al/Zr joining, including hardness and shear strength, were measured. The microhardness variation across the interface of steel/Al/Zr at temperatures of 545 °C, 560 °C, 575 °C and 590 °C are shown in Fig. 16. The value of the hardness in the reactive zones is higher than the matrix because of the formation of intermetallics (Al5Fe2 and Al3Zr). The average hardness values of Q345 steel, Al, Zr, Al5Fe2 and Al3Zr phases are 115 HV, 33 HV, 120 HV, 800 HV and 576 HV, respectively. The basic criterion for the formation of intermetallic phases in the diffusion zone is that the phase must first nucleate and then grow. The stoichiometry of the intermetallic phase may vary to some extents in various formation stage, influencing the hardness (Akhter et al., 2005). When the temperature is low at 545 °C, the diffusion reaction is not sufficient and Al5Fe2 phase layer is still in the first stage of formation. Therefore, the hardness at 545 °C is lower than those at higher temperatures. Similar phenomena have also been observed in FeNi3, σFeCrNi and Fe2Zr by Sabetghadama et al. (2010) and Aboudi et al. (2017), respectively. The variation trends of hardness along the interface are well identified with the EDS results, however, the hardness of second phases (Al3Fe and Al3Zr2) are not detected due to the very thin thickness of reaction layer. Fig. 17 shows the shear strength variation of the bonding at 545 °C, 560 °C and 575 °C. It is obvious that the diffusion temperature has a great influence on the interface shear strength. When the diffusion temperature is relatively low (530 °C), the bonding quality of interface between steel and aluminum foil is very weak due to the low thermal
3.6. Growth orientation law of Al3Zr phase The phase distribution in Al-Al3Zr-Zr region is shown in Fig. 12(a), where red represents Zr, yellow represents Al3Zr, blue represents Al, and white region represents unrecognized Al3Zr2. The grain sizes of Al and Zr are much larger than that of Al3Zr. For Al3Zr, the grain size decreases from Zr side to Al side. The size of Al3Zr mainly distributes from 0.9 to 3 μm shown in Fig. 12(b). The grain orientation distributions across the interface along RD, TD, and ND directions are shown in Fig. 13. Unlike Al5Fe2, there are not preferred orientations in Al3Zr along RD, TD, and ND directions. This is also confirmed from the full and inverse pole figures (Fig.14). In Al/Zr diffusion bonding, Kidson and Miller have revealed Kirkendall void formation at the reaction layer interface, which has also been observed in Fig. 2 in our study, and they observed an Al vacancy concentration gradient across Al3Zr (Kidson and Miller, 1964). It indicates that the growth of Al3Zr is controlled by the diffusion mechanism of aluminum vacancy (Fisher et al., 2013). Fig. 15(a) shows the grain boundary distribution of Al, Al3Zr and Zr. The pink lines stand for small-angle grain boundaries, and the black
Fig. 12. (a) The phase distribution of Al-Al3Zr-Zr region, (b) grain size distribution of Al3Zr. 8
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Fig. 13. The grain orientation distribution of Al-Al3Zr-Zr, (a) RD, (b) TD, (c) ND direction.
excitation of the atoms, and the shear strength is zero. When the diffusion temperature exceeds a certain value (590 °C), large crack forms between Al5Fe2 and aluminum foil owing to the higher volume fraction of brittle intermetallics, leading to the shear strength of zero. Therefore, only the shear strength at 545–575 °C are shown in Fig. 17. The shear strength has a small difference between 545 °C and 560 °C due to the similar thickness of reaction layer (545 °C: 83.6 um; 560 °C: 93.3 um). The thickness of reaction layer increases to 153.7 um at 575 °C, and the corresponding shear strength increases to 30 Mpa with a displacement of 0.16 mm. Above all, controlling a suitable temperature is very critical for diffusion bonding. A low temperature can not provide enough energy for diffusion reaction, while a high temperature will lead a drastic
reaction forming crack and voids. In order to analyze the shear fracture mechanism, the side and longitudinal fracture morphologies are shown in Fig. 18. From the side fracture morphology, one can see that the fracture of pure zirconium/ aluminum foil/Q345 steel all occurred in Al5Fe2 during the shear test from 545 °C to 575 °C, but exhibited a clear difference. At 545 °C, the thickness of Al5Fe2 layer was relatively thin, and a small amount of steel matrix material was also observed in the fracture (Fig. 18(a)). However, the fracture completely occurred within Al5Fe2 at 560 °C and 575 °C. The former was uneven, and showed obvious tearing trace, but the latter was flat (Fig. 18(c) and Fig. 18(e)). The longitudinal fracture at 545 °C showed that Al5Fe2 had a large plane-shaped morphology with
Fig. 14. The full (a) and inverse (b) pole figures of Al3Zr grains. 9
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Fig. 15. (a) and (b) are grain boundaries distribution of Al3Zr.
many gullies, which was caused by the mixed matrix material (Fig. 18(b)). As the temperature raised, the gully-like shape disappeared, and the plane-shaped Al5Fe2 splited into a rugged mass-like structure (Fig. 18(d)). When the temperature reaches 575 °C, the planeshaped Al5Fe2 transformed into a small, irregular blocky structure. There is no dimple shape, and the brittle cleavage fracture surface is shown in Fig. 18(f). The fracture mechanism is different to our previous study using Cu interlayer due to the different compositions of interface, in which the fracture is mixed with brittle and toughness in nature (Meng et al., 2018). The above microhardness analysis shows that the hardness of Al5Fe2 is about 250 Hv higher than Al3Zr, and the thickness of Al5Fe2 is more than three times of Al3Zr. Therefore, the fracture is more likely to occur in the region of Al5Fe2 during shear test. When the thickness of Al5Fe2 is moderate, the crack spreads completely within Al5Fe2 and the interface bonding strength is higher. When the temperature rises to 590 °C, the thickness reaches 191.2 μm, but it will cause a large internal stress due to the thick layers obtained at high temperature and short time, resulting in direct cracking between Al5Fe2 and the aluminum foil. Therefore, controlling a suitable temperature and extending the diffusion time can avoid the formation of big residual stress, and obtain a good bonding strength.
Fig. 17. Shear strength curves of joints at 545 °C, 560 °C and 575 °C.
Fig. 16. Hardness variation across the steel/Al/Zr interface at (a) 545 °C, (b) 560 °C, (c) 575 °C and (d) 590 °C. 10
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Fig. 18. The shear fracture morphologies of (a) side fracture at 545 °C, (b) longitudinal fracture at 545 °C, (c)side fracture at 560 °C, (d) longitudinal fracture at 560 °C, (e) side fracture at 575 °C, (f) longitudinal fracture at 575 °C.
4. Conclusions
Acknowledgements
The effect of temperature on the diffusion bonding of Q345 steel to zirconium using an aluminum interlayer was investigated. The reaction layers are mainly composed of relatively thick Al5Fe2 and Al3Zr phases with thin Al3Fe and Al3Zr2 as transition phases. Both Al/Fe and Al/Zr reaction layer thicknesses are proportional to the bonding temperature. The activation energy for Al3Zr is much higher than Al5Fe2, indicating that the formation of Al5Fe2 is easier than Al3Zr. The growth of Al5Fe2 and Al3Zr phases are controlled by the diffusion of Al and both of them preferentially form at Al side. Al5Fe2 has an obvious orientation with [001] ǁ TD, [100] and [010] ǁ RD and ND directions. The orientation of Al3Zr is totally random. The optimized shear strength can be obtained by choosing a suitable diffusion temperature, which determines the reaction degree and bonding quality between the matrix and Al interlayer. The hardness Al5Fe2 is much higher than Al3Zr, and the fracture always occurs within Al5Fe2 reaction layer.
This work was supported by the Fundamental Research Funds for the Central Universities (HEUCFP201850), Natural Science Foundation of Heilongjiang (LH2019E030 and JC2017012), the NSFC Funding (11874327, 51371062, 51701051, 51501192 and U1460102), China Postdoctoral Science Foundation Funded Project (2017M620111), and Innovation Center of Nuclear Materials for National Defense Industry (HCL-08). References Abbasi, M., Dehghani, M., Guim, H.U., Kim, D.I., 2016. Investigation of Fe-rich fragments in aluminum-steel friction stir welds via simultaneous Transmission Kikuchi Diffraction and EDS. Acta Mater. 117, 262–269. Aboudi, D., Lebaili, S., Taouinet, M., Zollinger, J., 2017. Microstructure evolution of diffusion welded 304L/Zircaloy 4 with copper interlayer. Mater. Design 116, 386–394. Ahmad, M., Akhter, J.I., Zaman, Q., Shaikh, M.A., Akhtar, M., Iqbal, M., Ahmed, E., 2003. Diffusion bonding of stainless steel to Zircaloy-4 in the presence of a Ta intermediate layer. J. Nucl. Mater. 317, 212–216. Akhter, J.I., Ahmad, M., Iqbal, M., Akhtar, M., Shaikh, M.A., 2005. Formation of dendritic structure in the diffusion zone of the bonded Zircaloy-4 and stainless steel 316L in the
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Journal of Materials Processing Tech. journal homepage: www.elsevier.com/locate/jmatprotec
Diffusion bonding of Q345 steel to zirconium using an aluminum interlayer a
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Yang Zhang , Bangzuan Long , Kai Meng , Aleksandr Gohkman , Ye Cui , Zhongwu Zhang
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Key Laboratory of Superlight Materials and Surface Technology, Ministry of Education, College of Materials Science and Chemical Engineering, Harbin Engineering University, Harbin, 150001, China Department of Physics, South Ukrainian National Pedagogical University, 65020, Odessa, Ukraine
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ARTICLE INFO
ABSTRACT
Associate Editors: C.H. Caceres and S.J. Na
The microstructure, interface reactions and mechanical properties of the Q345 steel/Al/Zr bonding interface were investigated. The thicknesses of the reaction layer in both Al/Fe and Al/Zr interfaces are temperaturedependent, with a higher temperature yielding a larger thickness. Both experimental results and thermodynamic calculation confirmed that the diffusion transition regions near the Al/Fe and Al/Zr interfaces mainly consist of Al5Fe2 and Al3Zr with thin Al3Fe and Al3Zr2 transitional phases, respectively. The growth kinetic equations of Al5Fe2 and Al3Zr as a function of temperature were investigated. The fitted curves from the equations well agree with the experimental results. The hardness of Al5Fe2 is higher than Al3Zr and the base materials. The propagation of shear crack occurs within Al5Fe2. Optimized shear strength can be obtained by controlling the thickness of reaction layer. The grain growth and fracture mechanisms of Al5Fe2 and Al3Zr at the interfaces were also discussed based on electron back scattering diffraction (EBSD) analysis of both reaction layers and fracture morphology.
Keywords: Diffusion bonding Microstructural evolution Growth mechanism Mechanical properties Reaction layer
1. Introduction Srikanth et al. (2017) reported that zirconium alloys are widely applied as the in-core structural material in various reactor vessels due to the superior properties. Whereas, stainless steels are Srikanth et al.usually used as the out-of-core components because of the good corrosion resistance. Therefore, bonding between steel and Zr-based alloys are often performed at various thermal reactors. However, Shaaban et al. (1978) found that direct joining of Zr-based alloys and stainless steel has encountered severe problems due to the different physical and chemical properties, such as the melting temperature and thermal expansion coefficient, which generate residual stresses leading to the microcracks formation at the bonding interface. Oliveira et al. (2016) indicated recently that the formation of brittle intermetallics (ZrFe2, ZrFe3 and ZrFe4) together with the crystallographic mismatch between Zr and Fe can reduce the bonding strength and lead to brittle failure. Oliveira et al. (2019) used friction element welding to keep the maximum welding temperature below the liquidus temperature of the high strength steel and friction element, which lower the risk in the formation of brittle intermetallic compounds. Ahmad et al. (2003) used electron beam welding to join Zircaloy-4 and stainless steel, although defects like porosity, voids and cracks were avoided, the harmful intermetallic compound Zr(Cr,Fe)2 formed. Zhang et al. (2019)
⁎
investigated the interfacial microstructure between Mg and 22MnB5 boron steel using resistance spot welding, and found that the source of crack was due to the formation of brittle Mg-Al phase. Bhanumurthy et al. (2001) found that the diffusion reaction was extremely sluggish between 403 stainless steel and Zr-2.5 wt% Nb alloy, and diffusion bonding could be conducted only at a high pressure and temperature (10 Mpa, 900 °C). Therefore, studies on the bonding of steel and zirconium should be performed to minimize or prevent the brittle phases formation to obtain a good bonding. A simple and feasible method for this aim is using an interlayer. Akhter at al. (2005) proposed that the interlayer can prevent brittle phases formation at the interface, and is helpful for reducing the bonding temperature and pressure. Previous studies used Cu, Ni, Ti, and Ta as interlayers to avoid the generation of brittle intermetallics between Fe and Zr. For example, Aboudi et al. (2017) found that Cu interlayer failed to restrict the formation of brittle phases but was benefit to reduce the hardness of the reaction region. Meng et al. (2018) studied the thickness effect of Cu interlayer on the microstructure evolution, shearing and bending properties. The optimized shearing and bending strength can be obtained by controlling a suitable Cu interlayer thickness. Srikanth et al. (2017) used Ni and Ti interlayers to join 304 L stainless steel to Zircaloy-4 alloy, and found that when temperature reached 850 °C and 900 °C, the joints would fail owing to the cracks
Corresponding author. E-mail address: [email protected] (Z. Zhang).
https://doi.org/10.1016/j.jmatprotec.2019.116352 Received 11 April 2019; Received in revised form 25 July 2019; Accepted 28 July 2019 Available online 30 July 2019 0924-0136/ © 2019 Elsevier B.V. All rights reserved.
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formation and propagation through the reaction layer (RL). Various methods have been conducted to join steel and zirconium, such as solid-solid diffusion bonding, explosion welding, fusion welding, and electron beam welded joints. These methods like explosion welding, however, have drawbacks of either high cost or lacking versatility. Perona et al. (1966) revealed that brittle intermetallic compounds, such as ZrFe2, ZrFe3 and Zr(Cr,Fe)2, were formed in the welded zone when using fusion welding joined stainless steel and Zircaloy-4These compounds are harmful to the mechanical property and corrosion resistance of the welded joints. However, many disadvantages of welded joints can be overcome using vacuum diffusion bonding process. Compared to the welding process, Springer et al. (2011) revealed that the base materials were not melted in diffusion bonding, which could avoid the residual stress of interface region in the joints caused by thermocycling. In the present study, the aluminum interlayer is used to bond Q345 steel and pure zirconium using diffusion bonding process. Kang and Kim (2015) have conducted on the joining of Al/Fe and Al/Zr. Their results confirmed that the diffusion bonding can provide a good strength. The choice of the base materials in this study are mainly due to the following reasons:
et al. (2007) found that the surface roughness had a large effect on the bonding, and thus all the base materials were polished carefully before bonding. Sand papers with various roughness of 800, 100, 1500, 2000, and 3000 were successively used to polish to a final roughness of about 1 μm. The polished base materials were cleaned by water first and further by absolute alcohol for 15 min in Ultrasonic instruments, followed by blowing dry with air blower. Diffusion bonding was performed from 530 °C to 620 °C with a pressure of 3 Mpa for 3 h in a vacuum hot press sintering furnace. The vacuum of the furnace is less than 1 × 10−3 MPa. A more detailed sample information and experimental procedure can be found in our previous study (Meng et al., 2018). The bonded samples at different temperatures were cut to sheets with a section of 12 × 12 mm2, and then the bonding surfaces were polished using mechanical polishing for microstructure observation and hardness measurement. At least five indentations (along the interface direction) were measured for each distance to obtain an average value. The microstructures were observed by metallographic microscope (LEICA DMIRM, Germany) and an FIE-Quanta 200FEG scanning electron microscope (SEM) with back scattered mode (BSE). The acceleration voltage is 20 kV. The phase compositions of the joint interfaces were identified by line scanning of energy dispersive spectrometer (EDS, FIE, Quanta 200FEG). X-ray diffraction (XRD) was performed at Rigaku D/max-TTR- III. To identify the products by XRD in the interface, the bonded samples were stripped layer by layer along the bonding interface to make sure the reaction products were exposed thoroughly. The reaction products at Zr and Fe sides were determined by XRD, separately. The microhardness of the bonding interface were conducted by an HVS-1000Z tester with 10 g load for 25 s. Shear test was performed using a San-Si tester. The schematic diagram of shear test can be found in our previous paper (Meng et al., 2018). A compression speed of 1 MPa/s was used. According to Yan et al. (2010), the shear stress (τ) was calculated from the following relations:
1 Q345 steel belongs to low-carbon high-strength steel which also has good plasticity and welding performance. The content of alloy elements is relatively low, which can avoid the effects of other alloy elements on the diffusion reaction between Fe and Al. 2 Pure zirconium has very low impurity content, and is easy for the mechanism investigation of diffusion bonding between Zr and Al. 3 Aluminum is chosen as an interlayer due to its relatively low melting point (660 °C) over other metals and, therefore, critical conditions for diffusion bonding can easily be met. Although many studies have been conducted on the joining of Al/ Fe, the interdiffusion and growth mechanism of interface remain contradictory. The most common view is that a large number of vacancies facilitate the fast diffusion of Al atoms, leading to the growth of Al2Fe5 along [001] direction (c axis) (Takata et al., 2015). However, the joining between Al and Zr using diffusion bonding is rarely investigated. Only interdiffusion and reaction were reported but the growth mechanism of interface remained unknown. Therefore, the diffusion bonding between steel and zirconium using an aluminum interlayer is of great importance for further understanding the interface growth mechanism. Mehrer (2007) found that diffusion in solids strongly depends on temperature, being slow at low temperatures but appreciable at high temperatures. Pressure also plays an important role, but is far less striking than temperature. Travessa et al. (2002) revealed that chemical reactivity at the interface is also a significant requirement for bonding, as it enables the occurrence of adhesion at reasonably low temperatures and moderate external pressures. The experimental conditions in this study can be identified as low temperature and low pressure incomparison to the common ones. The bonding at high temperature is easy to cause hydrogen absorption and lead to hydrogen embrittlement. Srikanth et al. (2017) pointed out that the thermal mismatch and intermixing can be limited by decreasing the bonding temperature. In this paper, effects of temperature on the microstructure, microchemistry of the interface, kinetic and thermodynamic features, hardness and shear properties are investigated in detail. The fracture mechanism, growth mechanism of the reaction phases at the interfaces are also discussed based on the fracture morphology and EBSD analysis of both reaction layers.
=
P ae
(1)
Where a and e are the width and height of test interface, respectively. In this study, a and e are 25 mm and 3 mm. The growth orientation law of Al-Fe and Al-Zr interfaces were studied by electron backscatter diffraction, EBSD (Hitachi S3400 scanning electron microscope equipped with HKL Channel 5 software). Fig. 1 shows the schematic diagram of coordinate system (RD-TD-ND) in EBSD measurements. 3. Results and discussion 3.1. Microstructure evolution Fig. 2 shows the interfacial microstructure evolution between 530 °C and 620 °C. The thickness of Fe/Al diffusion layer is almost zero at
2. Experimental procedure Pieces of Q345 steel (100 × 100 × 6 mm3) and pure zirconium (100 × 100 × 1.5 mm3) were used for vacuum diffusion bonding. A 500 μm thick pure aluminum foil was selected as an interlayer. Chen
Fig. 1. Schematic diagram of EBSD test coordinate system. 2
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Fig. 2. Microstructure of the bonding interface at (a) 530 ℃, (b) 545 ℃, (c) 560 ℃, (d) 575 ℃, (e) 590 ℃, and (f) 620 ℃.
530 °C, indicating that almost no reaction occur at this temperature. The thicknesses of both Fe/Al and Zr/Al reaction layers (RLs) increase as temperature increases, but with different changes. The increase in bonding temperature facilitates the atoms to diffuse across the interface, yielding a thicker reaction layer. The RL width of Al-Fe is almost 3 times wider than that of Al-Zr at 560 °C, indicating a higher reaction speed of Al-Fe than Al-Zr. From Fig. 2(a) to (f), one can observe that both interfaces develop towards interlayer Al owing to a larger diffusion coefficient of Fe and Zr to Al than that of Al to Fe and Zr at the same temperature (Fisher et al., 2013). When the bonding temperature reaches 590 °C, micro-crack emerges in the Al-Fe RL region. According to the binary phase diagram of Al-Fe, the Al-Fe compounds are all hard and brittle phases. In addition, Wang et al. (2014) reported that the thermal expansion coefficient of Al (23.5 × 10−6 K-1), Fe (12.2 × 10−6 K-1), and Al-Fe intermetallic compounds (18.94 × 10−6 K-1 for Al5Fe2, 19.68 × 10−6 K-1 for Al3Fe) are quite different. Thus, the microcracks are observed within Al-Fe intermetallic compound layer. With the further increase of temperature to 620 °C, almost all the Al interlayers have reacted with Fe and Zr, and large cracks emerge between the products of Al-Fe and Al-Zr. The thickness of Al-Zr products is larger than that of Al-Fe, indicating that bonding temperature has a totally different effect on the growth rate. Kidson and Miller (1964) also observed many voids due to the diffusion imbalance across the interface due to Kirkendall effect. The voids are near the Zr
side due to the higher diffusion coefficient of Zr to Al than Al to Zr. 3.2. Phase identification To analyze the elemental compositions across the reaction interface at 545 °C, 560 °C and 575 °C, EDS line scanning was performed and shown in Fig. 3. The element composition distributions are shown in Table 1. According to the atom ratio of Al/Fe and Al/Zr on the surfaces, the reaction products (Al5Fe2 and Al3Zr) were identified. Al5Fe2 exhibit a tongue-like morphology with highly irregular peaks and valleys orientated towards the iron. In the reaction of Fe and Al, Li et al. (2016) found that Al5Fe2 is always easily formed whether in solid phase reaction or solid-liquid reaction, and exhibit a morphology of wave shape towards the Fe direction. Dickson et al. (2014) found that thick Al3Zr reaction layers formed in the diffusion bonding of Al and Zr in the temperature range from 425 to 625 °C, and dense tiny voids were observed on the side close to Zr due to Kirkendall effect. Although this effect also exist in 545–575 °C, it was not observed because the atom diffusion was relatively slow in this temperature range. Only when the temperature reaches 620 °C, large voids form within Al3Zr due to the aggregation of vacancies as shown in Fig. 2f. In many binary diffusion couples, in addition to the formation of a stable phase, there will be one or more transition phases between the stable phase and the diffusion couple. The transition phases formed at 3
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Fig. 3. The chemical compositions across the bonding interface of steel/Al/Zr at (a) 545 ℃, (b) 560 ℃, (c) 575 ℃.
3.3. Phase formation sequence in Fe/Al and Zr/Al binary solid phase diffusion systems
Table 1 Elemental compositions of Al-Fe and Al-Zr reaction productions. T(oC)
545 560 575
Q345 steel side
Zr side
Fe(at.%)
Al(at.%)
Al/Fe
Phase
Al(at.%)
Zr(at.%)
Al/Zr
Phase
27.0 27.3 27.8
73.0 72.7 72.2
2.7 2.7 2.6
Al5Fe2
73.9 73.7 74.0
26.1 26.3 26.0
2.8 2.8 2.8
Al3Zr
The driving force of phase formation is mainly determined by the change of Gibbs free energy ΔG0, which can be expressed as
G0 = H 0
(2)
T S0
Where T is Kelvin temperature, and are the change in enthalpy (heat) and entropy, respectively. In solid-state reaction, Lee et al. (2003) revealed that the entropy change is further smaller than the enthalpy change, and thus it can be negligible. Then G 0 can be approximated by H 0 , indicating that the phase formation can be predicted by the heat of formation. Pretorius (1990) proposed that the formation law of products during solid phase diffusion can be well predicted according to EHF (effective heat of formation) theory. Pretorius et al. (1993) state that many factors such as impurities, lowest eutectic, atomic mobility, diffusing species, could affect the actual concentrations that are available for interaction at the growth interface. Therefore, the term effective concentration is used in the effective heat of formation concept. The released heat is dictated by the effective concentration of the limiting element (Pretorius et al., 1991). Therefore, the effective heat of formation can be described as follows:
H0
the junction of Al5Fe2, Al3Zr and the parent metal or interlayer. A high magnification SEM image of the RLs at 560 °C is shown in Fig. 4. The micron transition phases can be observed at the surfaces of Al5Fe2/Al and Al3Zr/Zr. The major elemental compositions at each spot within the transition phases detected by EDS show that the transition phases can be identified as Al3Fe and Al3Zr2, which has also been observed by Chen et al. (2016). The average thickness of transition phases also increase with the increase in temperature. The layered (Al3Fe) and tongue-like (Al3Zr2) morphology become more and more clear with the increase in temperature. The thickness of Al3Fe layer is much thinner than Al5Fe2 layer, which can be attributed to the much lower diffusion coefficient (DAl5Fe2 = 1.84 × 10−8 m2/s and DAl3Fe = 9.2 × 10-10 m2/s) (Bouché et al., 1998). The phase components of Fe/Al and Zr/Al interfaces were further verified by XRD (Fig. 5). From the patterns, it is found that the reaction product mainly consists of Al5Fe2 at steel side, and Al3Zr is observed on the zirconium side, which is well agreed with the EDS results and the following thermodynamic calculation results. However, the secondary transitional products, Al3Fe and Al3Zr2, are not detected due to the detection limit of XRD technology. The individual small unrecognized peaks maybe due to the oxide formation of Fe2O3 and Al2O3.
H = H0 ×
S0
effective concentration limiting element compound concentration limiting element
(3)
Theron et al. (1996) showed the formation enthalpy of the binary compounds in Al-Fe and Al-Zr systems. The lowest liquidus (eutectic) temperature is 660 °C according to the Al-Fe binary phase diagram, and the effective concentrations of Al and Fe at the interface are 99.1% and 0.9%, respectively. Thus, Fe is a limiting element. The effective heat of formation of different Al-Fe phases are shown in Fig. 6a, and it is obvious that ΔHʹAl5Fe2 < ΔHʹAl3Fe < ΔHʹAl2Fe < ΔHʹAlFe3. Al5Fe2 has the lowest effective enthalpy of formation, and Al3Fe takes second place, indicating Al5Fe2 and Al3Fe are the first and the second products, respectively, which further confirms the experimental results. According
Fig. 4. A high magnification SEM image of the RLs, (a) Al5Fe2/Al interface and (b) Al3Zr/Zr interface. 4
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Fig. 5. XRD patterns of (a) Fe/Al and (b) Zr/Al interfaces.
to EHF model, after the formation of the first product in binary metal system, the second product will form at the interface between the first product and the excess metal (M) and enrich the M element, and Pretorius (1990) confirmed this phenomenon in the experiments. Similarly, the lowest temperature on the liquidus curve is 660 °C in the Al-Zr binary system, and the effective concentrations of Al and Zr at the interface are 98.0% and 2.0%, respectively. In this case, Zr is considered as a limiting element. The effective enthalpy of formation of Al3Zr, Al2Zr, Al3Zr2 and AlZr were calculated and shown in Fig. 6(b). It is found that ΔHʹAl3Zr < ΔHʹAl2Zr < ΔHʹAl3Zr2 < ΔHʹAlZr, where Al3Zr has the lowest effective formation enthalpy of -3.28 kJ·(mol·at.)−1 being the first reaction product, which agrees with the experimental results. However, the calculated second product is Al2Zr rather than Al3Zr2 (experimental result). This phenomenon is named as “Phase jump behavior”, and has also been observed in Al-Pd (Colgan, 1987), Al-Pt (Hultgren et al., 1973) and Al-Au (Campisano et al., 1975) binary system. Laik et al. (2004) calculated the Gibbs free energy of Al2Zr or Al3Zr2 on the interface of α-Zr(Al)/Al3Zr by the solid solution model of binary system. The results show that ΔGAl3Zr2=-10.77 kJ/mol and ΔGAl2Zr=-7.18 kJ/mol, indicating that Al3Zr2 has a larger driving force and is preferentially formed between α-Zr(Al)/Al3Zr.
Table 2 The average thicknesses of the RLs. Temperature (oC)
Al5Fe2 (μm)
Al3Zr (μm)
545 560 575 590
84 93 154 191
20 31 65 124
diffusion temperature. Based on the experimental data at temperatures of 545 °C, 560 °C and 575 °C in Table 2, the relationships between lnx and T for Al5Fe2 and Al3Zr can be fitted as:
lnx = 9.22
15.25 ×
lnx = 25.78
30.00
1000 (Al55Fe2) T
(5a) (5b)
Then, the growth kinetic equations of Al5Fe2 and Al3Zr can be described as
xAl55Fe2 = 1.01 × 10 4 exp
126800 , QAl55Fe2 = 126.8kJ•mol RT
249400 , QAl3Zr = 249.4kJ •mol RT
1
(6a)
3.4. Growth dynamics of Al5Fe2 and Al3Zr
xAl3Zr = 1.57 × 1011 exp
The measured average thickness of the RLs from SEM-BSE are shown in Table 2. Luo and Acoff (2004) discovered the thickness of reaction layer, x, can be expressed as
The fitted and experimental temperature dependences of the Al5Fe2 and Al3Zr layer thickness are plotted in Fig. 7. The thicknesses of RLs from the calculations are well identified with the experimental results (the experimental result of Al5Fe2 thickness is zero and thus not shown here), which can be observed from the inset of Fig. 7. With the further increase in temperature, the growth curve of Al5Fe2 and Al3Zr intersects, and then the growth rates of these two phases begin to be reversal. From Fig. 2(e) and (f), one can observe that the thickness of Al5Fe2 is larger than Al3Zr layer at 590 °C, however, the thickness
lnx = ln (K 0 t )
Q RT
(4) −1/2
where K0 is a temperature dependent reaction constant (m·s ), t is diffusion time (s), R is Boltzmann constant (8.314 J·mol-1 K-1), and T is
Fig. 6. The effective heat of formation diagram, (a) Fe-Al and (b) Zr-Al. 5
1
(6b)
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interface reaction process and the growth rate is controlled by the diffusion of the slowest element. As the diffusion activation energy of Fe atoms in aluminum was less than that of Al atoms in the steel like mentioned above, and the Al-Fe interface grows towards to Al interlayer shown in Fig. 2, thus the growth of the Al5Fe2 phase can be assumed to be controlled by diffusion of Al. Abbasi et al. (2016) observed the presence of Al3Fe phases along the grain boundaries of Al, which offer fast diffusion paths for Fe atoms to reach low-Fe areas and nucleate. However, Al3Fe phase was not observed from the EBSD results due to the thin thickness. Fig. 9 shows the grain orientation distribution across the interface of RD, TD, and ND directions. Here, we mainly focus on the orientation distribution of Al5Fe2, for which red represents [001] orientation, blue represents [100] orientation and green represents [010] orientation. It is obvious that the [001] orientation of most Al5Fe2 grains is along TD direction from Fig. 9(b). The grains parallel to RD and ND directions have two types from Fig. 9(a) and (c): the blue ones with [100] orientation and the green ones with [010] orientation. To further investigate the orientation distribution of Al5Fe2 grains, the full and inverse pole figures are measured (Fig.10a and 10b). Ding et al. (2018) and Takata et al. (2015) used the similar method to illustrate the orientation of η-Fe2Al5 phase. It can be seen from Fig.10 that [001] concentrated along TD direction. The elongated grain growth of Al5Fe2 has a strong orientation, that is, the c-axis [001] direction is perpendicular to the interface between steel and aluminum, which is also confirmed by Murakami et al. (2004). This kind of grow priority guarantees that the Al5Fe2 phase grows much faster than other intermetallic phases. The grain boundaries distribution of Fe-Al5Fe2-Al system are shown in Fig. 11(a) and (b). In general, Onuki et al. (2013) and Zhang et al. (2018) defined that the grain boundary angle of 2°-15° belongs to a small-angle grain boundary (green line) and a grain boundary greater than 15° is defined as a large-angle grain boundary (black line). The focus is the distribution of grain boundaries within Al5Fe2, in which the large-angle grain boundary accounts for 70%, mainly distributed near the Al side; the small-angle grain boundary accounts for 30%, mainly distributed near the Fe side. Small angle grain boundary density is the highest in the area contacting with Fe, which suggests that the strain due to the transformation from α-Fe phase to Al5Fe2 introduces a high density of dislocations into the neighboring α-Fe phase, developing a dislocation substructure around the interface (Takata et al., 2015). The large Al grains mainly consist of small-angle grain boundaries, while the grains within Fe are mainly large-angle grain boundaries. Therefore, there is a grain boundary angle difference at the junction of Fe to Al5Fe2 transition, and the small angle grain boundary position has a higher dislocation density. Burkhardt et al. (1994) reported Al5Fe2 is an orthorhombic crystal structure with a space group of Cmcm (a = 0.766 nm, b = 0.642 nm, and c = 0.422 nm). While α-Fe belongs to a body-centered cubic crystal structure with space group IM3M (a = 0.293 nm). Since the number of atoms in Al5Fe2 and α-Fe unit cells
Fig. 7. The fitted and experimental curves between the thickness of RLs and temperature. The inset shows a zoom from 790 K to 870 K.
reverses at 620 °C. The interface microstructure at 620 °C can be used to explain this phenomenon shown in Fig.2(f), from which we can observed that the Al intermediate layer is almost consumed by Fe and Zr, and the thickness of Al3Zr is far greater than Al5Fe2. When the temperature reaches 617.7 °C, the reaction rate flips up, resulting in the thickness of Al3Zr begins to be larger than Al5Fe2 (Fig. 7). The activation energy for Al3Zr is 249.4 kJ mol 1, which is much higher than Al5Fe2 (126.8 kJ mol 1), indicating that the reaction of Al5Fe2 is easier than Al3Zr. This can be reflected by the RLs thicknesses of Al5Fe2 and Al3Zr as shown in Table 2. 3.5. Growth orientation law of Al5Fe2 phase In order to investigate the growth orientation law of interface phases during diffusion bonding, EBSD was used to characterize the microscopic crystallographic characteristics of Al5Fe2 and Al3Zr. The phase distribution of Al-Al5Fe2-Fe region are displayed in Fig. 8(a), where red presents Al, yellow presents Al5Fe2 and blue stands for Fe. One can see that Al and Fe have larger crystal grains than Al5Fe2. The grain size distribution of Al5Fe2 is shown in Fig. 8(b), and exhibits an obvious trend: the smaller grains are mainly concentrated at the Al side; the larger and elongated grains are more close to the Fe side, which agree with the previous study (Abbasi et al., 2016). It can be inferred that the crystal nucleation occurs at the Al side, and the fine crystal grains are mutually swallowed and gradually grow towards the Fe side with the increase of diffusion time. Laik et al. (2013) showed that only when the stoichiometric composition of phase layer is satisfied, the phase layer grows, and its growth process behaves as a series of
Fig. 8. (a) The phase distribution of Al-Al5Fe2-Fe region, (b) grain size distribution of Al5Fe2. 6
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Fig. 9. The grain orientation distribution of Al-Al5Fe2-Fe, (a) RD, (b) TD, (c) ND.
Fig. 10. The full (a) and inverse (b) pole figures of Al5Fe2 grains.
are different, leading to an apparent volume expansion originating from the transformation from the α-Fe to Al5Fe2 (Takata et al., 2015). Richards et al. (1994) points out that there is a large difference in thermal expansion coefficient, leading a difference in the stress at the transition interface of Fe to Al5Fe2. The Al5Fe2 grains suffer from stress during grain growth, causing that the grain boundaries perpendicular to the stress are stretched and those parallel to the stress are compressed. That is, the c-axis of the
grain growth direction is subjected to tensile stress, while a and b axis perpendicular to the grain growth direction are subjected to compressive stress. The grain of Al5Fe2 is elongated towards the c-axis, indicating that the stress in this direction is relatively small. Thus, the transformation from Fe to Al5Fe2 becomes easier. However, the stresses in a and b directions are large, leading to an obstruction of the growth direction in these directions. It is found that the crystalline defects of Al5Fe2 possess 30% vacancies along the c-axes ([001] direction), which 7
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Fig. 11. (a) and (b) are grain boundaries distribution of Fe-Al5Fe2-Al system.
offers a rapid diffusion path for Al5Fe2 and increases its growth rate, causing the formation of columnar Al5Fe2 and the corresponding residual steel substrate (Cheng and Wang, 2009). The growth of Al5Fe2 is controlled by the vacancy mechanism, and the grain boundary itself is the position of the vacancy source and the annihilation. The formation energy of the grain boundary vacancy perpendicular to the stress direction is low, and the number of vacancies is large; while the grain boundary vacancy formation energy of parallel stress is high, the number of vacancies is less, thus causing a certain vacancy concentration difference inside the grains. Therefore, a great quantity of vacancies gather in the c-axis direction, providing a favorable channel for the diffusion of atoms. As a result, the Al5Fe2 grains appear a slender morphology and the overall interface exhibits as a “wavy” shape.
lines represent high-angle grain boundaries. Then we mainly focus on the grain boundaries distribution of Al3Zr, and the misorientation angle distribution of Al3Zr is shown in Fig. 15(b). The large-angle grain boundaries possess 92.8%, and only a small part of small-angle grain boundaries are distributed in the grown up equiaxed grains. 3.7. Mechanical properties The mechanical properties of the steel/Al/Zr joining, including hardness and shear strength, were measured. The microhardness variation across the interface of steel/Al/Zr at temperatures of 545 °C, 560 °C, 575 °C and 590 °C are shown in Fig. 16. The value of the hardness in the reactive zones is higher than the matrix because of the formation of intermetallics (Al5Fe2 and Al3Zr). The average hardness values of Q345 steel, Al, Zr, Al5Fe2 and Al3Zr phases are 115 HV, 33 HV, 120 HV, 800 HV and 576 HV, respectively. The basic criterion for the formation of intermetallic phases in the diffusion zone is that the phase must first nucleate and then grow. The stoichiometry of the intermetallic phase may vary to some extents in various formation stage, influencing the hardness (Akhter et al., 2005). When the temperature is low at 545 °C, the diffusion reaction is not sufficient and Al5Fe2 phase layer is still in the first stage of formation. Therefore, the hardness at 545 °C is lower than those at higher temperatures. Similar phenomena have also been observed in FeNi3, σFeCrNi and Fe2Zr by Sabetghadama et al. (2010) and Aboudi et al. (2017), respectively. The variation trends of hardness along the interface are well identified with the EDS results, however, the hardness of second phases (Al3Fe and Al3Zr2) are not detected due to the very thin thickness of reaction layer. Fig. 17 shows the shear strength variation of the bonding at 545 °C, 560 °C and 575 °C. It is obvious that the diffusion temperature has a great influence on the interface shear strength. When the diffusion temperature is relatively low (530 °C), the bonding quality of interface between steel and aluminum foil is very weak due to the low thermal
3.6. Growth orientation law of Al3Zr phase The phase distribution in Al-Al3Zr-Zr region is shown in Fig. 12(a), where red represents Zr, yellow represents Al3Zr, blue represents Al, and white region represents unrecognized Al3Zr2. The grain sizes of Al and Zr are much larger than that of Al3Zr. For Al3Zr, the grain size decreases from Zr side to Al side. The size of Al3Zr mainly distributes from 0.9 to 3 μm shown in Fig. 12(b). The grain orientation distributions across the interface along RD, TD, and ND directions are shown in Fig. 13. Unlike Al5Fe2, there are not preferred orientations in Al3Zr along RD, TD, and ND directions. This is also confirmed from the full and inverse pole figures (Fig.14). In Al/Zr diffusion bonding, Kidson and Miller have revealed Kirkendall void formation at the reaction layer interface, which has also been observed in Fig. 2 in our study, and they observed an Al vacancy concentration gradient across Al3Zr (Kidson and Miller, 1964). It indicates that the growth of Al3Zr is controlled by the diffusion mechanism of aluminum vacancy (Fisher et al., 2013). Fig. 15(a) shows the grain boundary distribution of Al, Al3Zr and Zr. The pink lines stand for small-angle grain boundaries, and the black
Fig. 12. (a) The phase distribution of Al-Al3Zr-Zr region, (b) grain size distribution of Al3Zr. 8
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Fig. 13. The grain orientation distribution of Al-Al3Zr-Zr, (a) RD, (b) TD, (c) ND direction.
excitation of the atoms, and the shear strength is zero. When the diffusion temperature exceeds a certain value (590 °C), large crack forms between Al5Fe2 and aluminum foil owing to the higher volume fraction of brittle intermetallics, leading to the shear strength of zero. Therefore, only the shear strength at 545–575 °C are shown in Fig. 17. The shear strength has a small difference between 545 °C and 560 °C due to the similar thickness of reaction layer (545 °C: 83.6 um; 560 °C: 93.3 um). The thickness of reaction layer increases to 153.7 um at 575 °C, and the corresponding shear strength increases to 30 Mpa with a displacement of 0.16 mm. Above all, controlling a suitable temperature is very critical for diffusion bonding. A low temperature can not provide enough energy for diffusion reaction, while a high temperature will lead a drastic
reaction forming crack and voids. In order to analyze the shear fracture mechanism, the side and longitudinal fracture morphologies are shown in Fig. 18. From the side fracture morphology, one can see that the fracture of pure zirconium/ aluminum foil/Q345 steel all occurred in Al5Fe2 during the shear test from 545 °C to 575 °C, but exhibited a clear difference. At 545 °C, the thickness of Al5Fe2 layer was relatively thin, and a small amount of steel matrix material was also observed in the fracture (Fig. 18(a)). However, the fracture completely occurred within Al5Fe2 at 560 °C and 575 °C. The former was uneven, and showed obvious tearing trace, but the latter was flat (Fig. 18(c) and Fig. 18(e)). The longitudinal fracture at 545 °C showed that Al5Fe2 had a large plane-shaped morphology with
Fig. 14. The full (a) and inverse (b) pole figures of Al3Zr grains. 9
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Fig. 15. (a) and (b) are grain boundaries distribution of Al3Zr.
many gullies, which was caused by the mixed matrix material (Fig. 18(b)). As the temperature raised, the gully-like shape disappeared, and the plane-shaped Al5Fe2 splited into a rugged mass-like structure (Fig. 18(d)). When the temperature reaches 575 °C, the planeshaped Al5Fe2 transformed into a small, irregular blocky structure. There is no dimple shape, and the brittle cleavage fracture surface is shown in Fig. 18(f). The fracture mechanism is different to our previous study using Cu interlayer due to the different compositions of interface, in which the fracture is mixed with brittle and toughness in nature (Meng et al., 2018). The above microhardness analysis shows that the hardness of Al5Fe2 is about 250 Hv higher than Al3Zr, and the thickness of Al5Fe2 is more than three times of Al3Zr. Therefore, the fracture is more likely to occur in the region of Al5Fe2 during shear test. When the thickness of Al5Fe2 is moderate, the crack spreads completely within Al5Fe2 and the interface bonding strength is higher. When the temperature rises to 590 °C, the thickness reaches 191.2 μm, but it will cause a large internal stress due to the thick layers obtained at high temperature and short time, resulting in direct cracking between Al5Fe2 and the aluminum foil. Therefore, controlling a suitable temperature and extending the diffusion time can avoid the formation of big residual stress, and obtain a good bonding strength.
Fig. 17. Shear strength curves of joints at 545 °C, 560 °C and 575 °C.
Fig. 16. Hardness variation across the steel/Al/Zr interface at (a) 545 °C, (b) 560 °C, (c) 575 °C and (d) 590 °C. 10
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Fig. 18. The shear fracture morphologies of (a) side fracture at 545 °C, (b) longitudinal fracture at 545 °C, (c)side fracture at 560 °C, (d) longitudinal fracture at 560 °C, (e) side fracture at 575 °C, (f) longitudinal fracture at 575 °C.
4. Conclusions
Acknowledgements
The effect of temperature on the diffusion bonding of Q345 steel to zirconium using an aluminum interlayer was investigated. The reaction layers are mainly composed of relatively thick Al5Fe2 and Al3Zr phases with thin Al3Fe and Al3Zr2 as transition phases. Both Al/Fe and Al/Zr reaction layer thicknesses are proportional to the bonding temperature. The activation energy for Al3Zr is much higher than Al5Fe2, indicating that the formation of Al5Fe2 is easier than Al3Zr. The growth of Al5Fe2 and Al3Zr phases are controlled by the diffusion of Al and both of them preferentially form at Al side. Al5Fe2 has an obvious orientation with [001] ǁ TD, [100] and [010] ǁ RD and ND directions. The orientation of Al3Zr is totally random. The optimized shear strength can be obtained by choosing a suitable diffusion temperature, which determines the reaction degree and bonding quality between the matrix and Al interlayer. The hardness Al5Fe2 is much higher than Al3Zr, and the fracture always occurs within Al5Fe2 reaction layer.
This work was supported by the Fundamental Research Funds for the Central Universities (HEUCFP201850), Natural Science Foundation of Heilongjiang (LH2019E030 and JC2017012), the NSFC Funding (11874327, 51371062, 51701051, 51501192 and U1460102), China Postdoctoral Science Foundation Funded Project (2017M620111), and Innovation Center of Nuclear Materials for National Defense Industry (HCL-08). References Abbasi, M., Dehghani, M., Guim, H.U., Kim, D.I., 2016. Investigation of Fe-rich fragments in aluminum-steel friction stir welds via simultaneous Transmission Kikuchi Diffraction and EDS. Acta Mater. 117, 262–269. Aboudi, D., Lebaili, S., Taouinet, M., Zollinger, J., 2017. Microstructure evolution of diffusion welded 304L/Zircaloy 4 with copper interlayer. Mater. Design 116, 386–394. Ahmad, M., Akhter, J.I., Zaman, Q., Shaikh, M.A., Akhtar, M., Iqbal, M., Ahmed, E., 2003. Diffusion bonding of stainless steel to Zircaloy-4 in the presence of a Ta intermediate layer. J. Nucl. Mater. 317, 212–216. Akhter, J.I., Ahmad, M., Iqbal, M., Akhtar, M., Shaikh, M.A., 2005. Formation of dendritic structure in the diffusion zone of the bonded Zircaloy-4 and stainless steel 316L in the
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