MIG Brazing

This article investigates the MIG brazing capability of coated steel sheets for automotive applications, focusing on MIG brazing as a viable alternative to MAG welding, especially for chassis components. The study aims to evaluate the impact of coatings on brazing ability, wettability, and arc stability. MIG brazing is increasingly being used for joining Body-in-White parts, especially when Resistance Spot Welding is unsuitable. This research investigates the potential of brazing coated steels, noting that MAG welding often leads to defects such as blowholes.

MIG Brazing Process

MIG brazing is similar to MAG welding but utilizes a low melting point filler wire, typically copper alloys, that does not melt the substrates. It primarily employs pure argon as shielding gas and a short-circuit transfer mode. The main advantages of MIG brazing over MAG welding include:

• Lower heat input due to the lower melting temperature.
• Reduced distortion, making it suitable for thin sheets used in automotive applications.
• Improved visual appearance and less coating degradation, leading to better corrosion resistance.

Figure 1: MIG Brazing process with representative cross sections from brazing and welding

Methodology

The study focuses on overlap joints without gaps, using CuAl8 and CuSi3 filler wires with a diameter of 1.0 mm, and employs a short-circuit waveform at a welding speed of 500 mm/min. The results are expressed in terms of brazing range, focusing on criteria such as wettability of different coatings and spatter formation.

Key Findings

• Wetting Studies: Two methods were used to study wettability—real MIG brazed samples and a wetting pilot. The results indicated that higher wire feed speeds lead to increased current and voltage, which enhances heat input and improves wettability despite a rise in molten metal volume.

Figure 2: Effect of wire feeding speed and heat input on wetting angle using 0.8 mm thick sheets and CuAl8 filler wire

• Influence of Coatings: The coatings had no significant effect on the wettability or arc stability. The behavior of Zinc-coated steels were like that of bare steel, with AlSi coating showing a wider brazing range due to its thicker and more refractory nature.

Figure 3: Brazing range and wetting angle for different coatings using 0.8 mm thick sheets and CuAl8 filler wire

• Spatter Formation: Spatters measuring between 0.1 to 0.3 mm occurred, with a notable increase in spatter rates observed when using zinc coatings. This effect was consistent across various coatings.

Conclusion

The study concludes that the coating does not influence the MIG brazing capability, as the brazing range, wettability, and arc stability remained consistent. The AlSi coating exhibited a broader brazing range. Switching between coatings does not require a change in brazing parameters, although an increase in spatter is expected when brazing coated sheets in comparison to bare steels. In summary, MIG brazing is validated as an effective method for joining coated steel sheets in automotive applications, providing advantages in heat input and corrosion resistance over traditional arc welding. Source J. Haouas, MIG brazing ability of coated steel sheets for automotive applications, IIW 2020 conference, SC XVII  

Case Studies in Alloy Selection

• B-Pillar (Center Pillar)
• A-Pillar (Front Pillar)
• Exposed Quality Dual Phase Steel
• Align Alloy Selection to the Needs of the Part
  • Alloy Selection of 980 MPa Steel
  • Alloy Selection of 1180 MPa Steel

Automotive engineers have substantially more steel grades to select from in their quest to balance properties, performance, manufacturability, sustainability, and cost.

Conventional dual phase steels, with a microstructure of simply ferrite and martensite, have excellent formability in the drawing and stretching deformation modes. However, the characteristics of this phase combination that work very well in these deformation modes lead to challenges in bending and edge-stretch deformation.

Complex phase steels have superior performance in bending and edge-stretch deformation, but are not as good as comparable-strength dual phase steel.

3^rd Generation Advanced High Strength Steels are a family of grades that can combine the best features of other steels while minimizing some of their associated constraints.

Below are examples of how steel grades were applied to various parts in order to solve specific challenges.

B-Pillar (Center Pillar)

B-Pillars are a particularly challenging part for cold stamping applications.  The upper section must be of sufficiently high strength to prevent cabin intrusion during a side impact, while the lower section needs to maintain at least moderate ductility to absorb crash energy. Stiffness improves with deeper draw depth and more shape, yet the necessary formability to achieve these are typically limited by the high strength requirements.  Furthermore, the door opening regions require flanges to facilitate joining of outer and inner components.  Production manufacturing constraints dictate that the blank edges are usually formed by mechanical shearing rather than laser cutting, and forming the targeted part shape puts these cut edges in tension – exposing the risk for edge cracks of these higher strength steels. The combination of these challenges contributed to many OEMs choosing to form B-Pillars and entire door rings using hot stamping of press hardening steels.

Development of new types of advanced high strength steels bring the cold stamping option back into focus.

B-Pillar Upper

Metal flow when forming shear-cut blanks into B-Pillar Upper shapes puts the cut edges into tension along the front- and rear- door opening regions adjacent to the B-Pillar. Edge cracking propagating into the part, such as seen in Figure 1a, is frequently the outcome. Making this more challenging is that simulations have difficult predicting risk of cut edges, exemplified by Figure 1b which gives the false impression that there are no cracking or splitting concerns in this area.

Blank design countermeasures accomplish only so much. However, the steel industry now offers options at the same tensile strength but having better cut edge ductility as measured by the hold expansion test. The same part and process design is now capable of achieving the targeted part dimensions and characteristics, Figure 1c.

Figure 1: B-Pillar Upper stamped from conventional 590 GA (1a) and 590 GA with high hole expansion (1c). Figure 1b shows the simulation of the conventional 590 GA but does not indicate forming issues.^{S-125, J-30}

Table 1 presents a comparison of the properties used to form these parts.

B-Pillar Lower

OEMs usually design B-Pillar Lower portions to absorb side impact crash energy.  This means a combination of lower gauge and lower strength.  Vehicle stiffness must be maintained, so the lower horizontal portion may have a relatively deep draw depth – also needed to accommodate the contours of the rocker shape.  From a formability perspective, the deep draws and aggressive designs put these stampings (Figure 2a) at risk of splits due to insufficient elongation and n-value. An example of a split section is shown in Figure 2b. However, no issues were found when stamped from a 3^rd Generation 980 MPa steel, Figure 2c.

Figure 2: B-Pillar Lower (a) showing splits when made from conventional 980 GA (b), and split-free when made from 3rd Gen 980 MPa (c).^J-30  

A-Pillar (Front Pillar)

A-Pillar Lowers, also called front pillar lowers, have similar formability concerns as the B-Pillar Lower, except that A-Pillars are usually formed from even higher strength grades since they are at the front corners of the passenger safety cage.  The typical design leads to the fractures seen in Figure 3, arising from insufficient elongation and n-value in the steel.   Application of a 3^rd Generation 1180 MPa steel overcame these concerns, leading to robust and successful stamping.

Figure 3: Deep draw depths make A-Pillar Lower panels at risk for splitting. Converting to a 3rd Gen 1180 MPa steel may alleviate these concerns.^{S-125, J-31}

Application of Dual Phase Steels to Exposed-Quality Surface Panels

Bake hardening steels frequently are used for exposed panels due to the extra strengthening occurring from the paint curing step.  Of the bake hardenable grades, 340 BH is likely the most widely applied.  In this terminology typically used in some parts of Asia, these grades have a minimum tensile strength of 340 MPa, and a minimum yield strength of approximately 200 MPa.  As such, a 340 BH grade is similar to a 210 BH grade, using terminology more often associated with North America- and Europe based companies.

The option of a higher strength bake hardenable steel exists with a 440 DP product: a dual phase steel that has a minimum tensile strength of 440 MPa and has bake hardenability. The as-produced yield strength of 340 BH and 440 BH are similar, but the 440 BH grade has a larger strength increase associated with the paint curing step.^F-51 The low yield strength of 440 BH is attributed to having martensite in the ferrite matrix.  With a now higher strength panel, a thickness reduction of 0.05 mm is possible while still improving dent resistance.^J-29

Table 2 compares these steels, with YP’ representing the in-panel and baked strength, which is the sum of the yield point (YP), work hardening (WH), and bake hardening (BH).

Align Alloy Selection to the Needs of the Part

Alloy Selection of 980 MPa Steel

Body structures often call for parts to be formed from 980 MPa tensile strength steel.  Recognizing that strength is not the only characteristic that determines whether the part can be successfully and robustly stamped into the targeted shape and dimensions, the steel industry offers different Advanced High Strength Steel grades to meet the minimum tensile strength requirements.

Citations M-77 and M-78 compare three such grades: DP980, CP980, and QP980, all cold rolled steel at 1.6 mm thick.  The DP grade is a standard dual phase steel, CP represents a complex phase steel, and QP indicates that it is a Quenched and Partitioned 3^rd Generation AHSS grade.  Table 3 compares the tensile properties of the 3 steels, and Figure 4 shows the stress strain curves for these grades.  The bold portions of the curves in Figure 4b comes from tensile testing which generates data through uniform elongation.  From that point, the curves are extrapolated to the true fracture strain ε_f  determined by analyzing the dimensions of the fractured tensile bars. 

Figure 4: a) Engineering and b) True Stress-strain curves for DP980, QP980, and CP980.^M-77

Strain hardening was characterized as a function of strain, Figure 5.  The lack of a constant n-value prevents use of the Holloman relationship to characterize the hardening curves. Necking occurs once the true strain equals the work hardening coefficient (ε = n).

The initial peak in Dual Phase steel at low strains comes from the high work hardening offered by the microstructural ferrite, but decreases as the ferrite strengthens with increasing strain.  The plateau in the strain hardening coefficients for the QP980 and CP980 steels is consistent with the occurrence of the TRIP effect. The microstructural differences between the three AHSS types are directly related to their performance in formability characterization tests like tensile and bending.

Figure 5: Strain hardening coefficient as a function of strain for DP980, QP980, and CP980.^M-77

As indicated in Figure 6, the steels have different microstructures.  Dual phase steels are comprised of ferrite (α) and martensite (α’); the QP steel has ferrite (α), martensite (α’), and ≈6% retained austenite (γ); and the microstructural components of the CP steel include ferrite (α), martensite (α’), bainite (α_b), and ≈6% retained austenite (γ).

Figure 6: Microstructural components of DP980, QP980, and CP980.^M-77

In V-Bending tests described in Citations M-77 and M-78, the greatest strains were found in bends made on the DP980 steel, while the QP980 steel resulted in the lowest strains at the bend radius. The high n-value at high strains seen in QP980 reduces strain concentration along the bending line, and therefore reduces the peak strain at this location.  The springback angle was the largest in the CP980 steel, which is consistent with having the highest yield strength.  

Alloy Selection of 1180 MPa Steel

Stamped parts experience multiple modes of forming deformation, including drawing, stretching, and bending.  Some, but not all, advanced high strength steels having a multiphase microstructure are sensitive to cut-edge stretching as measured in a hole expansion test.  The key to stamping success is to deploy the correct steel grade that addresses the challenges and constraints of the part. AHSS grades offer options to address these challenges.

As one example, conventional 980DP and 1180DP was compared to 3^rd Gen 1180 in these various deformation modes.^M-54

The tensile and hole expansion properties of these grades are listed in Table 4, while Figures 7 to 9 shows the bendability, stretch formability, and deep drawability of the three grades.  The forming limit curve of the 3^rd Gen 1180 steel is comparable to 980DP. Additional testing confirms each of these steels are spot weldable with conventional techniques and have excellent resistance to delayed fracture associated with hydrogen embrittlement. ^M-54
 

Figure 7: Bendability of 3^rd Gen 1180 steel compared with 1180DP and 980DP. ^M-54

Figure 8: Stretch formability of 3^rd Gen 1180 steel compared with 1180DP and 980DP.^M-54

Figure 9: Deep drawability of 3^rd Gen 1180 steel compared with 1180DP and 980DP.^M-54

Grade Development Success Stories

• Galvanizing TRIP and 3rd Generation AHSS Grades
  
• New Alloy Product Development and Characterization: Complex Phase Steels

Galvanizing TRIP and 3rd Generation AHSS Grades

The background for this section comes from Citation B-93.

Galvanizing of TRIP and 3^rd Generation AHSS grades is particularly challenging. 

Manganese (Mn), aluminum (Al), and silicon (Si) are used in TRIP steel for strengthening, promoting austenite stability (Mn) and preventing cementite precipitation which indirectly enhances retained austenite stability (Si and Al).

The typical heat treatment used to produce the targeted TRIP microstructure in uncoated steels is not compatible with the hot dip galvanizing thermal cycle,  and results in a loss of the desired properties.

Heat treatment of uncoated TRIP steels usually starts with intercritical annealing (IA) typically at temperatures in the range of 750 °C to 900 °C, depending on the alloy and the targeted properties.  This is followed by holding at an isothermal bainitic transformation (IBT) of 360 °C to 410 °C to produce the best combinations of the proper microstructure, tensile strength, and uniform elongation.  However, this heat treatment cycle is not compatible with continuous galvanizing lines, where the zinc bath is typically held at around 465 °C and the incoming steel strip enters at a slightly higher temperature.^B-93

An IBT holding temperature at the higher temperatures typical of galvanizing undesirably impacts the microstructure developed at the lower IBT. The higher temperatures promote carbide precipitation, resulting in the final microstructure having a lower volume fraction of lower carbon retained austenite – neither of which are associated with good TRIP properties. This lower C content retained austenite will have lower stability, resulting in a rapid retained austenite to martensite transformation and mechanical properties characteristic of dual phase steels.^B-93

In addition, the speed at which coils move through continuous galvanizing lines dictates how long the steel spends in each step, including the galvanizing pot.  The time spent at the galvanizing pot temperature is shorter than necessary to produce the desired properties, at least with the chemistry used for uncoated TRIP steels.  TRIP steels with higher aluminum levels are potential alternatives for galvanized versions, since they may allow for targeted microstructure development at the shorter times associated with strip travel through the galvanizing pot, and result in adequate carbon enrichment in the retained austenite and as such providing the desired stability to this phase.  Furthermore, partially or completely replacing SI with Al has been shown to improve reactive wetting in galvanized TRIP steels.  Aluminum additions also may have the benefit of suppressing the depth and number of internal oxides, in particular at grain boundaries during coiling after hot rolling.^L-76

Furthermore, the primary alloying elements used in TRIP steel – Mn, Al, and Si – undergo selective oxidation during the annealing portion of continuous galvanizing lines.  Without proper control of the dew point and the oxygen potential of the annealing atmosphere, thick oxides cover the surface and result in poor reactive wetting and unacceptable bare spot defects in the Zn coating.^B-97  This applies to both hot dip galvanized^C-47 and hot dip galvannealed zinc coatings.^C-48

The presence of chromium (Cr) affects oxide layer formation, with the impact also being a function of the dew point and the oxygen potential of the annealing atmosphere.  Formation of a Cr₂O₃ oxide layer, which blocks penetration of additional oxygen to the steel beneath it, is the same mechanism that makes stainless steels “stain less” and improves corrosion resistance.^Z-22

Poor wettability is shown in Figure 1, using an example from Citation D-49.

Figure 1: Poor wettability after galvanizing. SiO2 covers the surface in the bare areas, while the sparse areas covered with zinc correspond to reaction areas with an interfacial Fe2Al5 layer in between the zinc and steel. Modified from Citation D-49.

The dew point below which wettability problems can start is a function of the alloy chemistry and other conditions in the annealing furnace.  Figure 2 highlights the transition from a non-wettable to wettable surface as the dew point increases from −35 °C to −26 °C.

Figure 2: Wettability improves as the dew point increases from −35 °C to −26 °C.^D-49

 Other Strategies to Improve Galvanizability

As reported in Citation G-57, methods used to improve the galvanizability of TRIP steels include:

• Use of a high dew point atmosphere, resulting in the internal selective oxidation of Si and Mn to SiO2 and xMnO.SiO2, respectively;
• Pre-deposition of a thin layer of Fe, Ni or Cu to prevent selective oxidation altogether.
• Pre-oxidizing the steel surface to form isolated, rather than continuous, oxide particles which are embedded in a thicker Fe-oxide surface layer. This layer is subsequently reduced to metallic Fe in a reduction step.

These approaches are summarized in Figure 3.

Figure 3: Potential methods to promote zinc wettability and avoid bare spots during hot dip galvanizing of TRIP steels.^G-57

TRIP steels contain several elements with strong tendencies to oxidize, namely manganese (Mn), and silicon (Si), in addition to the iron (Fe) in the alloy.  Carbon plays no direct role in surface oxidation but can be lost by decarburization at high dew points where carbon combines with oxygen to form either CO or CO₂.

By using a slightly oxidizing annealing atmosphere and a slightly higher dew point than typical annealing atmospheres (approaching 0°C, instead of −40 °C),^{B-94, B-95, B-96, N-34} it is possible to preferentially oxidize Fe to form a thin surface FeO layer before Mn and Si can oxidize externally. Manganese and silicon are prevented from reaching the surface, and any Mn or Si oxides form beneath the iron oxide layer or at grain boundaries rather than at the top surface. 

Next, the steel passes through a strongly reducing atmosphere at a lower dew point (−20 °C or preferably lower)^{J-32, G-57}  and greater hydrogen concentration. 

The iron oxide surface layer is reduced back to metallic Fe by hydrogen, producing iron and water vapor. The internal oxides (MnO, SiO₂) remain beneath the surface where they do not impact zinc wettability.

The steel surface after this oxidation-reduction process is primarily iron, offering substantially improved zinc wettability and adhesion with a reduced risk of bare spot defects in the galvanized zinc coating.

Deployment of this oxidation-reduction method relies on accurate control of atmospheric conditions at the oxidation and reduction chambers.

New Alloy Product Development and Characterization: Complex Phase Steels

In addition to being able to produce a new grade, it is important to consider the applications where it might be used, and the characteristics it must have.  In the past, it might have been sufficient to design only for tensile strength and possibly ductility as measured by total elongation in a tensile test.  But the industry is demanding much more than that from today’s advanced steels.

An example of some of the efforts involved in new alloy characterization are found in Citation R-30, where advanced testing involved with commercializing a hot rolled High Strength Low Alloy steel and a hot rolled Complex Phase steel are highlighted.

Metallurgically, the desire was to move away from single-phase ferritic or bainitic steels that rely on micron-sized second-phase constituents for higher strength, but also produce damage and voids upon shearing, to an approach that derives strength from nanometer-sized precipitates that do not promote substantial void formation during shearing.

The typical evaluation for cut edge stretchability is the hole expansion test, standardized under ISO 16630.  This test is used to assess stretch-flangability and to predict edge-cracking during manufacturing. The test evaluates the steel’s resistance to through-thickness crack propagation in a shear-affected and work-hardened zone from
punching.

However, the stress and strain conditions imposed on edges during high volume manufacturing may be more complex than those found in the hole expansion test.

The study in Citation R-30 used a deep-draw-cup (DDC) test to screen multiple alloying approaches of the targeted grades for crack susceptibility and fracture toughness.  This outcome was used further to study the relationship between hole expansion and toughness, and the associated contributions of microstructural features like grain size, second-phase fraction, and texture.  Additional information about the specifics of the tests and results are found elsewhere.^R-31

Alloy composition impacts strength and formability, but also influences corrosion resistance, which is an important consideration in hot rolled steels for frame and chassis applications.

For the alloying, different balances of niobium, vanadium, and titanium, and molybdenum were used to influence austenite recrystallization and r-values in the HSLA grades with the target of suppressing the formation of carbon-rich second-phase constituents, other than nanometer-sized carbide and carbonitride precipitates. In the complex phase steels, boron and titanium were used to promote a titanium carbide-strengthened single-phase bainitic microstructure.  In these, the C content was varied to promote differences in carbon-rich second-phase fractions in the bainitic microstructure (martensite and cementite).

The mill processing route impacts properties.  After heating to 1240 °C, lab samples were rolled down to 3.5 mm. Finish-rolling temperatures (FRT) of 980 and 890 °C were used to study the effect of finish rolling on planar anisotropy and grain size. After hot rolling, the plates directly transferred to a run-out-table (ROT) and cooled with a mixture of water and pressurized air before being placed in a furnace to replicate coil cooling.  For the HSLA steels, the exit ROT and furnace temperatures were 580 and 630 °C. For the complex phase steels, the exit ROT and furnace temperatures were 400 and 450 °C. The effect of a change in microstructure (e.g., grain size, second-phase constituents) on HEC and toughness was investigated by varying the coiling temperature (CT).

Microstructural characterization was done with Scanning Electron Microscopy (SEM) and Electron Backscatter Diffraction (EBSD), in addition to X-Ray Diffraction (XRD) for texture analysis.

Each condition was evaluated for tensile properties, as well as DDC testing for susceptibility to crack formation and the energy required for crack growth, Hole Expansion testing, and Charpy V-notch (CVN) impact testing at -60, -40, -20, 0, and +20 °C to determine the absorbed energy and ductile-to-brittle transition temperature.

Even though the hole expansion values between the HSLA approach and the CP steels were similar, the bainitic CP steels exhibit substantially lower crack susceptibility and increased toughness than ferritic HSLA steels in the deep-draw cup (DDC) test. The DDC test was capable of detecting differences in crack susceptibility and work of fracture for different steel types under test conditions that simulate industrial manufacturing conditions.

Concurrent with these microstructural and mechanical tests, samples were evaluated for corrosion resistance using SAE J2334, VDA 233-102, and VDA 621-415.

No one alloy chemistry and processing route resulted in a single answer having the best performance in all of the characteristics evaluated, but this wide scope of testing provided the data that the steelmaker used in commercializing a new grade, HR660Y760T-CP. 

Instead of the conventional focus on stretch-flangability and hole expansion, the approach was widened to additionally include fracture toughness to assess crack susceptibility under conditions that take into account secondary cold work embrittlement and replicate component manufacturing through successive forming operations with alternating compressive and tensile stresses.  The detailed investigation also provided guidance on which alloying elements could be reduced to result in more cost-effective production without compromising performance.