Abstract:
In this study, friction welding process was successfully applied to join a ferritic ductile iron. The microstructures, phase transformation, crack analysis, microhardness and tensile test of the welds were investigated. The welding region was composed of the deformed graphite nodule, coarse pearlite, proeutectoid ferrite and acicular martensite. Highly deformed graphite nodules were distributed along the weld interface due to the flow of the plasticised material in the thermo-mechanically affected zone (TMAZ). In the central zone, graphite displayed a striped configuration and ferrite transformed into martensite structure. In the peripheral region, graphite surrounded by martensite remained as individual granules. The microhardness test results show that the maximum hardness at the interface reached 605 HV. The microstructure analysis shows that the cracks initiate mostly at the interface of the deformed graphite nodules and then spread through the grain boundaries of metal matrix. It is shown that the cracking process is predominantly controlled by the residual stresses. The fracture surface appearance shows a cleavage fracture mode in the peripheral zone and a little dimple fracture mode around the graphite nodules in the central zone. Moreover, the microcracks were identified at the graphite-matrix interface.
Key Words: friction welding; ductile iron; microstructure; cracks; microhardness; fractography
1 Introduction
Ductile iron containing spheroidal graphite is much stronger and has a higher elongation than other grade cast irons. Relatively high strength and toughness of ductile iron are advantageous in terms of many structural applications. This high ductility allows the castings to be used in critical high stress applications such as: crankshafts, steering knuckles, brackets, valves, truck axles, hubs, water pipes and many others [1]. Mechanical properties of ductile irons are directly related to their matrix microstructure. As-cast matrix microstructure of ductile irons may be entirely ferritic, entirely pearlitic or a combination of ferrite and pearlite, with spheroidal graphite distributed in the matrix. These microstructural features are affected by the solidification-cooling rate associated with the section size of the castings as well as that of the alloying elements [2].
Many techniques and special materials are available for the repair welding of ductile iron castings, for joining ductile iron to itself or to other ferrous materials. These welding methods include shielded metal arc welding, flux cored arc welding, gas metal arc welding, gas tungsten arc welding, and submerged arc welding, diffusion bonding, impact-electric current discharge joining, laser welding, oxyacetylene powder welding and lately friction stir processes [3].
When ductile iron castings are being repaired or joined by traditional fusion welding technologies, their high carbon content (more than 3.5%) may cause the formation of hard brittle phases in the fusion zone (FZ) and heat affected zone (HAZ). The carbon atoms of ductile iron diffuse into austenite during welding and form martensite and carbides at the weld interface while cooling [4]. These give rise to poor elongation properties and high hardness values, as reported by Pascual et al. [5]. These structural transformations result in reduction of ductility to a level where susceptibility to cracking is so high that either spontaneous post-welding cracking of the joint takes place, or cracks are generated when the first operational load is applied [6]. Therefore welding ductile iron, likewise welding other cast irons, requires special precautions to obtain optimum properties in the welded metal and adjacent HAZ. Brittle martensite can be tempered to a lower strength, more ductile structure through preheating or postweld heat treatments such as stress-relief annealing. Some welding procedures are designed to reduce the size of the HAZ and thus minimize cracking [7].
While the main goal is to avoid the formation of excess cementite in the matrix material, which makes the welded region brittle, in ductile iron an additional objective concerning the retaining a nodular form of graphite, is of almost equal importance [8]. To minimize the formation of massive carbides and high-carbon martensite, the most helpful is to have carbon presents in the form of spheroids which have a low surface-to-low volume ratio [9].
Friction rotary welding (FRW) is an example of welding process in which both, similar and dissimilar metals can be joined by solid-state diffusion processes to overcome metallurgical complications associated with fusion welding. This welding method produces a weld when two or more workpieces, rotating or moving reciprocally, are brought into contact under pressure to produce heat and plastically displace material from the weld interface [10]. The friction welding process is beneficial in terms of reducing the cost of complex forgings or castings, for example, the welding of a spindle or shaft to a cast/forged head [11, 12]. However, in case of friction welding of steels a lot of defects occur, including: centre defects, restraint cracks, weld-interface carbides, hot-shortness cracks or porosity, which may lead to failure of the friction welds [13].
The main problem with the friction welding of ductile iron are free graphite precipitations which have lubricating properties that reduce the efficiency of the welding process. During the friction process, the graphite spheroids are deformed or fragmented, thus creating an unfavourable microstructure [14-16]. Consequently, all the properties relating to strength and ductility decrease as the proportion of nonnodular graphite increases, and properties relating to failure, such as tensile strength and fatigue strength, are more affected by small amounts of such graphite than properties not involving failure, such as proof strength. The form of nonnodular graphite is important, because thin flakes of graphite with sharp edges have more adverse effect on crack initiation [13].
Another problem is the absence of interatomic bonds in friction welding of ductile iron caused by forming of continuous graphite between contact surfaces. This is also accompanied by a large decrease of the friction coefficient, which in turn decreases the amount of heat generated at the interface. Higher pressure in rotation of the components should support the heating of the joint to the melting point. However, an increase of pressure is accompanied by the formation of the network of cracks in the weld zone due, to the small plasticity of ductile iron [17].
In recent years a lot of studies concerning continuous friction welding [18-21] and friction stir welding (FSW) processes [22-24] of ductile iron to ductile iron or ductile to other low carbon metals were conducted by various researchers. Knowledge available from friction weld processes works concentrated on the structural and mechanical properties [25, 26] and metallurgical phase transformation [27] or welding parameters optimisation [28, 29]. A crack and failure characteristics of ductile iron during friction welding process hitherto have not been reported in the literature. Only other materials, such as mild steel and medium carbon steel were considered in an application concerning friction welding of truck axle, where the spindles are welded to the axle housings [30].
The main purpose of the present research is to show the microstructure and cracks features observed in friction welding of ductile iron. Furthermore, we took a closer look into metallurgical phenomena, accompanying the friction welding of ductile iron. As such, the Vicker’s hardness test, tensile test, optical and electron microscopy techniques were used in this study.
2 Experimental
The chemical composition of base materials selected for the study is shown in Table 1. The microstructure of the as-cast ductile iron showed a bull’s eye structure with ferrite surrounding the graphite nodules in a pearlitic matrix (Fig. 1). The specimens in the shape of bars of 20 mm in diameter and 100 mm length were cut for friction welding. The surface for friction welding was prepared on the abrasive cut-off machine. The process of joining was carried out on the continuous drive friction machine of a ZT-13 type. Heat force (HF) and upsetting force (UF) used in the experiment were in the range of 20-40 kN. A friction time (FT) of 10-30 s was applied. The upsetting time (UT) used in the experiment equaled 3 sec. The spindle rotating speed (SRS) was kept constant at 1450 rpm. Room temperature tensile tests were carried out as per ASTM:E8/E8M-11[31] standard specimen configuration. The tensile strength test was carried out on a 100-kN servo-controlled universal testing machine (Instron). Moreover, Vicker’s microhardness measurements were made across the weld line with a load of 500 g and a hold time of 15 sec. The microstructure of joints was examined by means of light metallography (LM) as well as a scanning electron microscopy (SEM). The specimens were mechanically polished by using emery special sheets with the help of disc and bench polishing machine. The prepared specimens were etched by applying 3% nital for inspecting the metallurgical behavior of the welded joints. The fracture surfaces of the specimens were observed in the SEM using BEI COMPO mode, applying backscattered electrons (BSE).
3 Results
3.1 Macrostrucure observation
A typical ductile iron friction welded joint and cross-sectional macrographs are shown in Fig. 2.A flash formed on the ductile iron sides is given in Fig. 2(a). Different characteristic zones, such as weld interface, thermo-mechanically affected zone and HAZ are analysed in detail on the etched cross-section of the joint, shown in Fig. 2(b). The microstructure variations of different zones in ferritic ductile iron are discussed below.
The fracture appearances of the interface after the tensile test are shown in Fig. 3a-b. It can be clearly seen that the shining region near the periphery is not involved in the bonding. The incomplete bonding was detected at the center of the joint, as shown in Fig. 3a. The middle portion of the friction weldment is directly involved in the bonding and the fracture region may vary from smoother to rougher areas. The spiralling marks show that the bond is formed through the rotational friction under the influence of axial load.
According to AWS [11] the low rotational speed and axial load led to this incomplete bond. Optimal rotational speed and axial road will be able to generate sufficient frictional heat flux during the process which will lead to the plasticizing of the interface. In the case of base metal fracture specimens (Fig. 3b), no such unwelded interface regions can be observed.
3.2 Microstrucure observation
During FRW of ductile iron, metal within the HAZ of the weld is heated above the eutectoid transformation temperature A1 and transforms into austenite. At the end of the forging stage, this metal cools below the eutectoid temperature and austenite decomposes into ferrite, pearlite, bainite or martensite as daughter products [32, 33]. The results of the microstructure observations of the friction welded joints as function of distance from the interface were given in Fig. 4. As can be seen in Fig. 4, a thin proeutectoid ferrite (PF) layer forms close to the weld interface in both, central and peripheral zone. Similar phase transformation after friction welding of carbon steel was observed in the papers [34-35]. This result suggests that the process temperature in TMAZ is high enough to austenitise and coarsen the grains. A relatively high cooling rate results in the coarse pearlite (P) (see Fig. 5a) and proeutectoid ferrite (Fig. 4 and Fig. 5b) precipitating at the prior austenite grain boundaries [36]. Moreover, the ferrite structures in the original ductile iron transformed into acicular martensite (M) structures (Fig. 4 and Fig. 5c) through rapid cooling from the high temperature state [33]. When the temperature in the ductile iron exceeds eutectoid temperature, the carbon in graphite spreads out to speed up the microstructure transformation into extent by increasing the carbon concentration in the base metal. Whilst rapidly cooling, the austenite structure may transform into martensite structure [37-40].
The microstructures of ferritic ductile iron in the HAZ included irregular and deformed graphite (DG) precipitates and a mixture of pearlite and ferrite. The proportion of the pearlite was small in comparison with that of the ferrite and ferrite was mainly presented around the region of graphite nodules. The results of microscopic observations show that the graphite morphology (size and shape) was being changed within an increasing distance from the weld interface. In some areas of the micrograph, the ultra-refined graphite (UG) particles (0.5-1 µm) were uniformly distributed in a pearlitic matrix (see Fig.4). The spheroidal graphite morphology changed into almost flake-shaped (FG) graphites located along the weld interface because of the extensive mechanical deformation and heat produced during FRW [23, 24]. The length of FG was approximately 50 µm (see Fig. 4). Using a high magnification, as given in Fig. 5(d), numerous dense acicular martensite structures around the graphite nodules were distinctly presented. The martensite particles have approximately (2-6 µm) in size. Based on the phase transformation to martensite, the temperature of the HAZ during welding should reach the eutectoid temperature. According to Cheng et al [22], because the remaining time of eutectoid temperature was extremely short, only the structures around the graphite transformed into martensite while cooling.
As can be seen from Fig. 5(e) for ductile iron, the process of friction welding is difficult because the graphite particles, distributed in metal matrix without ductility, degrade the deformation plastic flow in high temperature and act as a lubricant not allowing for sufficient frictional heat to process the material. Because the plastic deformation of the spheroidal graphite during RFW was significant, the deformation layers of spheroidal graphite (LDSG) were observed (Fig. 5e-f). The increase in the amount of deformation in thickness (about 40 %) leads to elongation of graphite through the metal matrix. It is clear that, at this high amount of reduction, the matrix is distorted around graphite, especially at both sharp ends of angular graphite and in the vicinity of graphite clusters creating areas of striations. These striations pass through many grains without deviation resulting in severe localized micro-cracking which is often observed along the matrix grain boundaries. These micro-cracks are sufficient to produce incipient fractures as reported by El-Bitar et al. [41]
According to Mitelea et al. [14], where short friction time and high axial pressures are concerned, it is not possible to weld the two materials together because of an early softening of ductile iron and a constant transfer of graphite nodules to the joint plane, leading to the formation of a new fresh graphite film (see Fig. 5(f)). However, the appropriate welding conditions reduced the amount of the LDSG in the joint of ductile iron [18].
3.3 Microhardness
Figure. 6 shows the Vicker’s hardness profile of the ductile iron cross-section following FRW. The hardness of the TMAZ, HAZ and a parent material region were measured. The hardness value of the ductile iron parent material was approximately 185-195 HV. According to the microhardness curve of the peripheral zone, the maximum hardness value reached 605 HV in TMAZ because numerous martensite structures were observed in this region (see Fig. 5c). It is thought that this structure was generated because the material was intensively heated and rapidly cooled during the FRW process [23, 24]. High hardness of the peripheral zone suggests that the carbon content of the martensite structures was relatively high. The result shows the hardness of the welded specimen decreases more rapidly when the phase changes from fully deformed zone to heat-affected zone. As can be observed from the Fig. 6, the hardness decreases in the HAZ much slower and extends to 7 mm from the weld interface. The hardness value of this region ranged between 210 and 310 HV, which remained higher than that of the parent material. The hardness value of the central zone was slightly lower than that of the peripheral zone, and the hardness of the parent material was the lowest. The maximum hardness value of the central zone, indicated in TMAZ, reached 516 HV. As expected, the hardness indicated close to the weld zone was much higher than the hardness of the HAZ and the parent material. Changes of hardness in the welding interface are directly associated with the microstructure which resulted from the degree of the heat being introduced and plastic deformation[3]. Earlier studies indicated that when carbon alloy is used in FRW, refined grains are generated at the weld because of the dynamic recrystallisation caused by severe plastic deformation at high temperatures. The structures of these refined grains comprised a lot of dislocation phenomena as reported by authors[42, 43].
3.4 Tensile strength, flash diameter and axial shortening
The tensile test was applied after machining the weld flashes, formed during the friction welding process. The effect of the friction time, flash diameter and axial shortening on the tensile strength of welded joints has been presented in Fig. 7a-c. The friction time was changed from 10 to 30 s. The samples were welded by constant heat force and the upsetting axial force. Generally, the results of the tensile strength obtained for ductile iron specimens are not satisfactory. As can be observed from the diagram in Fig. 7a, the tensile strength of the joints increased with the increasing friction time for ductile iron samples. Similar results of the relationship between the matrix structure and tensile strength yet, instead, for dissimilar joints were achieved in papers[25, 26]. As illustrated in the diagram (Fig.7b) the flash diameters increase with the increasing friction time for ductile iron welded joints. The relationship between friction time and axial shortening are presented in Fig. 7c. It is clearly seen that axial shortening increased with increasing friction, heat force and friction time. The combination of high friction welding parameters produced more axial shortening that played an important role in the mechanical properties [25, 26].
3.5 Cracking Observation
Figures 8(a) and (b) show the observed cracks. It may be seen that there are multiple cracks and the deformed graphite nodules in the path of the cracks. The cracks are visible around the graphite and some cracks are actually in the graphite (circle in Fig. 8a). More investigation on a situation, when there was a lower cracks severity, showed that indeed the graphite deformed nodule particles are the initiation sites for the cracks (arrows in Fig. 8(a)). Therefore, when a crack initiates, some of the thermal stress is released. However, when in the adjacent areas the thermal stress is further developed, another crack initiates from the graphite particle and of course propagates in the metal matrix (Fig. 8(b)). The connection of these minor cracks will lead to the formation of large major cracks, which run through the martensite structures (see Fig. 5(c)). Similar crack mechanism has been observed in case of powder welding process in papers [44, 45].
Figure 9 shows a schematic representation of the cracking in the HAZ of a ductile iron in case of FRW method. As can be seen, the marked areas are critical, because the highest residual stress exists in these areas. The ellipsoidal or even flake-shaped graphite in the HAZ acts as a stress raiser, which may prematurely cause a localized plastic flow at low stresses and initiate fracture in the matrix at higher stresses (Fig. 9). Moreover, the graphite flakes deflect a passing crack and initiate countless new cracks as the material breaks. As a result, this graphite morphology exhibits no elastic behaviour and fails in tension without significant plastic deformation [13, 46]. The details of the fracture surface characteristics in ductile iron welds are described in the section below.
3.6 Fractography
The fracture surfaces of the tensile tested specimens were characterized by means of scanning electron microscopy so as to understand the failure patterns. The SEM micrograph of the fracture surface of the tensile specimen is presented in Fig. 10.
As can be seen from Fig. 10a-c, the two dissimilar fracture morphologies have been distinguished in the ductile iron welds. The fracture observation of the peripheral zone sample showed a cleavage fracture with the rivers markings on the facets (red arrows in Fig. 10(a). Rivers markings on the facets result from the propagation of the crack on a number of planes of different levels [47]. The cleavage planes {100} or {110} types are clearly observed in the immediate vicinity of graphite nodule. These cleavage planes are characteristic for ferrite, which forms a specific shell around graphite nodule [27]. The fracture observation in the central zone sample showed the ductile mode of the fracture surface (Fig. 10b). Moreover, a microvoid coalescence (MVC) seems to be the dominant form of this fracture region (Fig. 10c). Additionally, Fig. 10c shows the microcracks which formed at the graphite-matrix interface, surrounded the nodule particles and then propagate to the matrix. A similar observation was also reported by Askari et al. [48].
4 Conclusions
Friction welding process was successfully applied to ductile iron. The examinations and analyses of microstructures and microhardness were performed. Furthermore, crack analysis, graphite particles morphology and phase evolution were primarily addressed in this work. The following conclusions are drawn from this work:
(1) During FRW of the ductile iron both, the high temperature and carbon content, significantly affected microstructure transformation. The final microstructures in the HAZ of ductile iron weld were composed of pearlite, proeutectoid ferrite and acicular martensite around the graphite nodules.
(2) The graphite in the surface zone exhibited a striped configuration, and distinct martensite structures formed in the matrix. The graphite in the central zone remained as individual granules, and acicular martensite was observed outside the graphite nodules.
(3) The microstructure analysis shows that the cracks initiate mostly at the interface of deformed graphite nodules and then spread through the grain boundaries of the metal matrix. It is shown that the cracking process is predominantly controlled by the residual stresses.
(4) The ellipsoidal or even flake-shaped graphite in the HAZ acts as a stress raiser, which may prematurely cause a localised plastic flow at low stresses and initiate fracture in the matrix at higher stresses.
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