1. Introduction
Due to stricter environmental regulations being adopted worldwide in recent years, the automobile industry needs to develop technology for lightening a vehicle body as a preemptive action for fuel efficiency improvement in addition to the reduction of CO2 emission. In general, aluminum (Al) or magnesium (Mg) alloys are commonly used for lightening a vehicle body, but high-strength steel of 1.5 GPa or higher is increasingly being applied for specific parts such as A-pillar, center pillar, and roof rail which require hardness and durability as they are directly related to the safety of passengers. In particular, the thickness can be reduced compared to general steel materials by applying high-strength steel to core parts of collision preparation, while productivity can be expected to increase through the shortened process time and cost reduction by minimizing the use of reinforcing parts1-3). However, molding into complex shapes is limited due to insufficient ductility during press molding as steel materials gain high strength, and the mold life is shortened due to an increased load applied to molds. As a solution, the technology of applying hot stamping, which enables molding at high temperatures while ensuring hardness and moldability through quick cooling, to boron steel with excellent hardenability is becoming commercialized.
Hot stamping is a thermal process at high temperatures of 900℃ or above in which scale generation and surface decarburization may occur due to surface oxidation of steel materials if oxidation caused by atmospheric gas is not prevented4). Therefore, various coating processes such as hot dip galvanizing or Al-Si coating are applied in order to prevent oxidation on steel material surface before heating. However, the Al-10Si coating layer developed by ArcelorMittal is most frequently applied currently based on outstanding oxidation and corrosion properties4,6) since liquid metal embrittlement occurs in a galvanized steel plate as liquefied zinc diffuses to austenite grain boundaries, thus inducing coagulation, due to low solubility of zinc during the hot stamping process5). According to a study by Cheng et al.6), porosity and defects in the Al-Si coating layer decreased as the silicon (Si) content in Al-Si coating layer increased in the Ai-Si coating layer, and the thickness of an intermetallic compound layer was reduced and uniform coating properties were secured by up to approximately 10% Si content.
On the other hand, laser weldability has the advantage of reduced thermal deformation compared to other types of welding due to high-speed and low heat input welding, in addition to fine bead appearance, precision, and stability. Due to such characteristics, leading automobile manufacturers around the world have already introduced and utilized laser welding technology for welding a vehicle body, and domestic companies are also considering laser welding as the next-generation manufacturing technology for complementing the conventional welding system7). According to a study by Kim et al.8), however, the coating layer is mixed into the fusion zone(FZ) during laser welding of Al-Si coated boron steel which results in concentrated precipitation in certain areas without being uniformly distributed within the FZ and therefore, ferrite is formed which causes the hardness of the FZ to decrease. As a solution to the above problem, mechanical grinding, chemical etching, and laser ablation processes have been used to remove the coating layer in advance8-11). Furthermore, additional factors such as a filler wire have been used for segmentation of the FZ and controlling a ferrite phase fraction in order to stabilize the austenite phase. Research is actively conducted in recent years as diverse processes such as arc preprocessing or beam oscillation are being developed12-16).
This paper examines previous studies on the characteristics and major impacts of laser welding of Al-Si coated hot-stamped boron steel, and aims to provide data necessary for practical research and development by analyzing domestic and international R&D trends related to weldability improvement from various perspectives including material and process aspects.
2. Characteristics of Al-Si coated Hot-stamped Boron Steel
2.1 Characteristics and microstructure of hot-stamped steel plate
Since the hot stamping process is metamorphosis reinforcement from hot pressing and die-quenching, steel materials used in the process must have a chemical composition with improved hardenability to be transformed from austenite phase in hot temperature to martensite phase through an appropriate cooling speed. In general, hardenability is improved by adding alloying elements such as B, Mn, Mo, and Cr in which the composition of these elements is determined considering mechanical properties obtained after heat treatment and economic feasibility. Particularly, hardenability is substantially improved in steel with a minimum amount (0.001 - 0.003 wt%) of boron (B) added because boron is segregated to the austenite grain boundary, thus delaying the heterogeneous nucleation of ferrite17,18). Fig. 1 shows the microstructure of boron steel before and after hot stamping. As shown in Fig. 1(a), the boron steel plate consists of ferrite and a small amount of martensite positioned at the grain boundary between pearlite and ferrite, and transforms into almost 100% late martensite through hot stamping as shown in Fig. 1(b). According to a study by Farabi et al.20), tensile properties are reported to decrease as martensite of the heat affected zone softens due to tempering by welding heat. As a solution, tailor welded blanks (TWB) are currently used in which hot stamping is performed after laser welding7).
Fig. 1
Typical base metal microstructure of the investigated steels, (a) as-received condition, (b) hot-stamped condition (modified from ref.19))
2.2 Al-Si coating layer structure
Based on the Fe-Al binary phase diagram, Fe and Al are not mixed with each other to form five types of intermetallic compounds including Fe3Al, FeAl, FeAl2, Fe2Al5, and FeAl3 (Fig. 2). FeAl2, Fe2Al5, and FeAl3 phases with a high content of Al have high hardness and brittleness, whereas Fe3Al and FeAl with a high content of Fe show improved fracture toughness as well as excellent wear resistance, oxidation resistance, corrosion resistance, and specific strength properties22). These intermetallic compounds are controlled by austenitizing heat treatment time and temperature according to the diffusion principle in which the extent of diffusion of Fe tends to increase as heat treatment time lengthens or temperature increases.
The Al-Si coating layer, which performs the role of preventing oxidation and decarburization of the basic material during hot forming, typically has the thickness of approximately 25-30 µm, the chemical composition of 88.73% Al, 10.31% Si, and 1.96% Fe, and the melting point of 600℃6,23). Fig. 2 shows the SEM image of the Al-Si coated boron steel coating layer before and after hot stamping. Fig. 3(a) shows the coating layer before heat treatment in which the intermetallic compound layer of Fe2SiAl7 with the thickness of 4 - 6 µm was formed along with precipitates that have grown into the Al-Si matrix, while the intermetallic compound of Fe2Al5 and FeAl3 was thinly distributed in the region where the baseplate and the coating layer are in contact. In Fig. 3(b) on the other hand, the coating layer thickness increased to approximately 10 - 15 µm on average through the diffusion process between the baseplate and the Al-Si coating layer during heat treatment; Fe combines with Al through mutual diffusion from the basic material to the coating layer to form intermetallic compounds, and oxides are formed due to the oxidation reaction with atmosphere19,23). Furthermore, a complex alloy layer having a band form of an intermetallic compound of FeAl2 and Fe2Al5 is created as the Fe content within the coating layer increases. These Fe-Al-based compound layers have a higher melting point of 1,100 - 1,300℃, which is two times higher than that of the Al-Si coating layer, and higher electrical resistance and brittleness, thus impeding weldability4,23).
Fig. 3
Cross-sectional scanning electron microscope (SEM) micrograph of a type aluminized coating, (a) before hot stamping, (b) after hot stamping (modified from ref.23))
3. Laser Weldability Consideration
3.1 Laser weldability of Al-Si coated boron steel
Laser welding types are largely divided into the keyhole mode where beads are formed due to fusion and evaporation of molten metal, and the conduction mode where beads are formed as nearby solids are melted due to heat conduction of melting liquid. The components of the basic material and the coating layer are mixed through mutual mixing of liquid and coating layer in the fusion zone formed in the keyhole mode, whereas liquid is formed in the fusion boundary without mixing in the heat conduction mode24,25). Fig. 4 presents the mechanical properties of the FZ according to laser output. When a complete joint is already formed in the keyhole mode at 1.1 kW output, tensile strength and microhardness of the FZ decreased as the output increased because the cooling speed of the FZ decreased from the increased heat input and the grain size of the FZ was enlarged25,26). Therefore, it is important to form a stable weld zone through the keyhole mode, but process optimization is required to select the heat input appropriate for the applied parts considering grain size and cooling speed.
Fig. 4
Mechnical properties of welded joints at different laser powers; (a) Tensile strength, (b) microhardness distribution curve (modified from ref.26))
Fig. 5 shows the EPMA result and the cross-section of the laser weld of Al-Si coated boron steel. Lei et al.26) used a high-speed camera to observe that white metal vapor is formed as the Al-Si coating layer was partially vaporized during laser welding. As shown in Fig. 5, however, most of the coating layer was mixed in the fusion zone during laser welding, and concentrated precipitation occurred in certain areas without uniform distribution within the FZ, thus forming a secondary phase of a white band; the content of Al and Si tended to be high in order to improve the phase stability thermo- dynamically. In particular, the Al content was high locally with an average of around 1.7 - 2.5wt%. The microstructure of the laser weld of Al-Si coated boron steel is shown in Fig. 6 for additional analysis. Fig. 6(a) shows that most of the microstructure of the FZ is composed of martensitic elements, but the Fe-Al phase is distributed between certain epitaxy grains of a mesh form. Yoon et al.27) confirmed the Fe3(Al, Si) phase through the EDS analysis result and the Fe-Al-Si ternary phase diagram. In addition, the long δ-ferrite phase was formed along the fusion line of the laser weld as shown in Fig. 6(b). Compared to the martensite matrix formed throughout, Fe3(Al, Si) and δ-ferrite phases have a relative hardness value of 50 - 60%, thus being soft and easily causing plastic deformation. Therefore, the dislocation ring moves to the second phase region when the applied load reaches beyond a certain level, which results in the growth and propagation of microcracks through micropores, thus causing brittle fracture7,29). Consequently, the second phase from the mixing of the Al-Si coating layer formed in the laser weld was confirmed to consist of Fe-Al intermetallic compounds, which are in a brittle phase that degrade mechanical properties, and soft ferrites.
Fig. 5
(a) Optical Microscope image with a rectangle highlighted for elemental mapping with EPMA for (b) Al and (c) Si of laser welded cross section (modified from ref.27))
3.2 Ferrite phase change affecting mechanical properties of fusion zone
Al in the Fe-Al binary phase diagram is a ferrite stabilizer element that reduces the austenite region of a single phase while expanding the double phase of δ- ferrite and austenite13). Martin et al.30) reported that the formation of the austenite phase is completely suppressed if the Al content exceeds 1.2wt%. Fig. 7 shows the Fe-Al binary phase diagram based on thermodynamics in which the effects of Al content on phase transformation characteristics and δ-ferrite phase formation are demonstrated. According to the phase diagram, a peritectic reaction (δ + L → δ + γ) completely transforms the liquid in which the δ-ferrite phase, or the high-temperature stable phase, is first formed in the liquid in a molten pool during cooling. Ultimately, the final microstructure has a double phase of ferrite and martensite if the Al content of the FZ exceeds approximately 0.3wt% during a phase transformation into austenite. In other words, the Al-Si coating layer mixed due to high welding heat during laser welding stabilizes the δ-ferrite phase in a molten pool. Then, δ-ferrite which is a non-equilibrium phase is maintained along the austenite grain boundary due to laser welding properties of quick cooling, thus suppressing a phase transformation from δ-ferrite to austenite28).
Fig. 7
ThermoCalc binary phase diagrams of Fe-Al and illustrated influence of Al on the phase transformation and δ-ferrite formation (modified from ref.19))
Fig. 8 presents the schematic diagram of phase changes in the basic material and the FZ during the TWB process in which hot stamping is performed after laser welding. The weld zone after laser welding has mostly consisted of martensite, but δ-ferrite is formed along the fusion boundary due to the Marangoni effect and the mixing of the Al-Si coating layer. δ-ferrite phase exhibits irregular patterns because nucleation and growth of martensite during laser welding have accelerated, which resulted in greater stress by nearby martensite with enlarged particle size31). Ferrite is diffused and formed throughout the FZ of the Al component mixed in during hot stamping heat treatment, and stabilizes into the double-phase structure of martensite and α-ferrite after die-quenching. As a result, stress is removed during austenitizing heat treatment of the δ- ferrite phase under a high compressive stress, and then the δ-ferrite phase is stabilized to the α-ferrite phase which further reduces the martensite fraction of the FZ29,31). Recently, Kang et al.30) conducted a study on the effects of heat treatment temperature and time on the microstructure characteristics of the laser welding. The Al content of ferrite decreased as aluminum diffused into the basic material when the austenitizing temperature and time increased. However, such process optimization cannot completely improve the mechanical properties from the mixing of the Al-Si coating layer, and there are extensive restrictions on practical applications at production sites considering economic costs and the characteristics required for the applied parts.
Fig. 8
Schematic of mechanism for fusion zone microstructure evolution; (a) after laser welding, (b), (c) annealing up to austenitizing temperature, (d) after the die quenching (modified from ref.13))
4. Laser Weldability Improvement Technology
4.1 Laser improvement technology in the material aspect
During laser welding and hot stamping of Al-Si coated boron steel, a brittle fracture may occur at the weld zone due to the double-phase structure of ferrite and martensite resulting from the mixing of Al in the FZ. Therefore, a complete austenitizing treatment is mandatory in order to prevent the ferrite phase from forming throughout the FZ for the prevention of brittle fracture12,28).
First, the most fundamental method is to apply mechanical grinding and chemical etching to the Al-Si coating layer on the surface of the laser FZ to remove the layer in advance in order to prevent the mixing of the coating layer in the FZ. The microstructure obtained from the SEM analysis of the TWB and the laser weld depending on the presence of the Al-Si coating layer is shown in Fig. 9. Overall, the FZ has mostly consisted of martensite. Fig. 9(a) and (c) show the δ- ferrite phase resulting from the mixing of Al-Si coating layer in the FZ, and then it transformed into the α-ferrite phase, or the equilibrium phase, through hot stamping heat treatment. In contrast, the laser weld for which the coating layer has been removed mainly consists of complete lath martensite as shown in Fig. 9(b) and (d) in which the lath martensite becomes finer after the hot stamping process. The reason is that martensite was transformed from columnar dendritic austenite particles after laser welding, which was then recrystallized during hot stamping heat treatment after die-quenching to undergo a phase transformation from finer equiaxed particles of austenite13,31). This method has the advantage of being able to easily improve the properties of the FZ at a low processing cost, but it requires an extensive amount of time and has low reproducibility, thus being difficult to apply in actual mass production.
Secondly, when carbon steel filler wire is applied, strong austenite stabilizing elements (C, Mn, etc.) are added to the FZ, the microstructure of the FZ will be fully composed of austenite at high temperature and martensite after die-quenching to room temperature. Based on the thermodynamic calculation, Fig. 10 shows the increase in carbon content from 0.2wt% to 0.6wt% results in the microstructure of the FZ undergoing a complete phase transformation into austenite at 950℃, thus forming martensite after die-quenching to room temperature. Specifically, the average Al content and the δ-ferrite volume fraction in the FZ both decreased when carbon steel filler wire was added. The δ- ferrite distribution within the FZ was more uniform when a filler wire was used; δ-ferrite was less and finely distributed because mixing was more active in the region with a broad flow of a molten pool as the feeding speed of a wire increased or the welding speed decreased12). Consequently, when the Al content is less than 3.7wt%, the most effective method is to use a filler wire containing strong austenite stabilizing elements during laser welding. Tensile strength and elongation of the FZ without a filler wire were approximately 1,210 MPa and 1.1%, respectively, while those of the FZ with a filler wire were around 1,546 MPa and 5.5%, thus resulting in improved physical properties and no brittle fracture.
Fig. 9
SEM microstructure in fusion zones before hot stamping, (a) coated welded joint, (b) de-coated welded joint and after hot stamping, (c) coated welded joint, (d) de-coated welded joint (modified from ref.9))
Fig. 10
Effect of carbon on Fe-Al binary phase diagram for Fe (C, Mn, Si, Cr)-x%Al of Al-Si coated boron steel grade in this study calculated using Thermo-Calc™ (modified from ref.13))
In addition, Khan et al.33,34) recently conducted a study on improving weldability by additional coating with graphite and nickel (Ni) which are austenite stabilizing elements on top of the existing Al-Si coating layer. The microstructure of the TWB according to the presence of an additional graphite coating layer is shown in Fig. 11. Fig. 11(a) and (d) show that the basic material is identical under the two conditions with respect to the fusion boundary, but the microstructure of the FZ varied significantly. Fig. 11(b) and (c) show that the particle sizes of ferrite and martensite are almost identical, but Fig. 11(e) and (f) in which a graphite coating layer is added show mostly martensite excluding a few small ferrite particles (< 3 μm) found in the FZ and along the fusion boundary. When a graphite coating layer with a thickness of 130 μm is used, the ferrite fraction of the FZ drastically decreased from 40% to 2%, which enabled the implementation of desired mechanical properties and fracture mode. Since carbon, which is an austenite stabilizing element, plays the role of a physical barrier as well as insulation capturing the Al-Si coating layer, a greater amount of the Al-Si coating layer is discharged to the outside through vaporization due to welding heat generated from a laser26). On the other hand, the nickel coating layer sufficiently secured the martensite fraction of the FZ even at the thickness of around 50 μm34). This result implies that the mechanical properties of Al-Si coated boron steel can be improved through laser welding of materials added with a filler wire and additional coating layer even without removing the coating layer in advance.
Fig. 11
SEM images of the base metal (BM), fusion boundary (FB), and fusion zone (FZ) in TWB joints (a) no additional coating condition with (b), (c) showing magnified images and (d) 130 μm graphite coating condition with (e), (f) showing magnified images (M: Martensite, F: Ferrite) (modified from ref.33))
4.2 Laser improvement technology in the process aspect
The attempts to actualize effective laser welding have led to the development of various laser welding processes as the applications of Al-Si coated hot-stamped boron steel has expanded. Laser ablation creates the same effect as the process of removing the Al-Si coating layer through chemical etching and mechanical grinding explained above. This process involves removing the Al-Si coating layer by using short multiple impulse lasers before laser welding where the process can be quantitatively controlled, unlike mechanical grinding or chemical etching, only leaving the Fe-Al intermetallic compound at the bottommost layer. Moon et al.10) discovered that mechanical properties can be improved by preventing the ferrite formation in the FZ only when laser ablation was applied to both sides of the coating layer; tensile properties and hardness improvements were not achieved when laser ablation was applied to only one side for cost reduction. Therefore, additional process design is needed for applying laser ablation to both sides, and weak corrosion resistance around the FZ was concerned if the coating layer was removed during the manufacturing process.
Arc preprocessing can also be applied without removing the Al-Si coating layer before laser welding. Fig. 12 shows the SEM and ESD analysis results to compare the characteristics of the Al segregation vulnerable part according to the application of arc preprocessing. The Al-Si coating layer was transformed into a complete intermetallic compound layer through arc preprocessing, and the Al content was uniformly distributed throughout at the fusion boundary. Therefore, a uniform hardness distribution in the FZ was achieved by increasing the martensite fraction in the FZ and decreasing Al segregation through arc preprocessing, and fracture occurred in the basic material during the tensile strength test.
Fig. 12
SEM images and EDX mapping for Al and Si element distribution after laser welding (modified from ref.14))
Dual beam laser welding technology using lead beam and lag beam has recently been developed as well as spatial modulation technologies such as laser beam oscillation. Fig. 13 shows the schematic diagram of each welding technology. For dual beam laser welding shown in Fig. 13(a) and (b), welding in the keyhole mode was stabilized and the fluid flow in a molten pool was improved according to a study by Li et al.15). Fig. 14 shows that dual beam laser welding provides a broader heating region and a longer molten pool than single beam laser welding based on the high-speed camera results. Therefore, segregation can be reduced as the Al content becomes more uniformly distributed within the FZ as the molten pool stabilizes into a broader region. However, when the Al content in the FZ is uniformly distributed due to dual beam laser welding, mechanical properties may be degraded as more regions reach the threshold of the ferrite phase region depending on the thickness and characteristics of the coating layer11,16). Accordingly, dual beam laser welding technology requires a complex optimization process by applying along with a filler wire and additional coating layers mentioned in Section 4.1.
Fig. 14
Evolution of the keyhole and molten pool during (a-c) single beam laser welding and (d-f) dual beam laser welding (modified from ref.36))
Laser beam oscillation technology illustrated in Fig. 13(c) can increase the convection of a molten pool by inducing nonlinear motions in the keyhole mode. Furthermore, a longer time is required for a molten pool to completely coagulate as input heat lasts and a molten pool is broader than laser linear welding. Fig. 15 shows the Al content distribution within the FZ of laser linear welding and laser beam oscillation welding. For laser linear welding, the Al concentration was higher around the fusion boundary and underneath the FZ compared to other regions. Laser beam oscillation welding, on the other hand, generates an efficient agitation effect, which reduced Al segregation through strong convection in a molten pool, while Al distribution became more uniform as the lasing diameter increased. However, the macro-segregation of Al concentrated in the fracture-vulnerable area of the FZ was mitigated through laser beam oscillation technology, but the ferrite phase formed within the FZ could not be completely removed; thus, a complex method of welding an additional thin nickel plate was applied.
Fig. 15
Concentration maps for solutes Al in the weld joints: (a) laser linear welding, (b) laser oscillation welding with a 0.5 mm diameter, (c) laser oscillation welding with a 1.3 mm diameter (modified from ref.16))
Ultimately, weldability can definitely be improved by solving the problem of ferrite formation within the FZ during laser welding of Al-Si coated boron steel by using diverse methods from optimization of laser welding process to adjusting the alloy composition of the FZ in a complex manner owing to recent technological breakthroughs; however, further research is needed from process optimization to commercialization.
5. Conclusion
This paper introduced the overall research results related to mechanical properties and microstructure of Al-Si coated boron steel applied with laser welding. The current research trends on materials and processes for improving laser weldability were examined from various perspectives, and the following conclusions have been obtained.
1) Regarding the mechanical properties during laser welding and hot stamping of Al-Si coated boron steel, the strength of the weld zone is decreased and brittle fracture may occur along the fusion boundary due to the formation of ferrite or intermetallic compounds within the FZ from mixing and diffusion of Al.
2) By primarily removing the Al-Si coating layer before laser welding, the quality of the FZ can be enhanced by preventing Al segregation and ferrite formation within the FZ. Coating layers can be removed through mechanical grinding, chemical etching, or laser ablation; improvements in mechanical properties can be achieved only if both sides of the coating layer are removed since welding in the keyhole mode is primarily performed.
3) The method for applying additional coating layers and a filler wire containing strong austenite stabilizing elements were studied in order to avoid brittle fracture in the laser weld through a complete austenitizing treatment throughout the FZ. In particular, the δ-ferrite distribution was more uniform than when a filler wire was used based on a phase diagram and microstructure during laser welding applied with a carbon steel filler wire, while tensile strength and elongation were improved. In addition, additional coating layers drastically reduced the ferrite fraction within the FZ by playing the role of a physical barrier and insulation that encloses the Al-Si coating layer.
4) When various technological trends such as arc preprocessing, dual beam, and beam oscillation for improving laser weldability were examined, uniform Al distribution was achieved and macro-segregation in the FZ could be prevented by stabilizing welding in the keyhole mode and improving the fluid flow in a molten pool. However, materialistic elements such as a filler wire and additional coating layers may be applied in a complex manner in order to completely remove the ferrite phase formed within the FZ.
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