A Novel Joint of 18-8 Stainless Steel and Aluminum by Partial Welding Process to Ni-Plated Carbon Fiber Junction

Abstract

A novel joint of 18-8 stainless steel and aluminum (Al) connected by a 20 × 10 × 0.3 mm nickel (Ni)-coated carbon fiber (CF) cross-weave cloth junction utilizing extremely large friction force by broad interface of 6 µm diameter CF has been successfully developed for potential aerospace application. This was done by a partial welding process of the CF junction to 18-8 stainless steel and Al separately taking advantage of the different welding temperatures of 660°C for Al, and about 1400 to 1450°C for 18-8 stainless steel. First, the 18-8 side is fabricated by inserting half the CF junction length (10 mm) into a cut slit in 18-8 rod half length, followed by the welding and rapid solidification. Next, the remaining exposed CF junction half length is inserted into Al rod, also followed by welding and rapid solidification, resulting in the finished [18-8/CF/Al] joint. The partial welding was performed by novel method of spot electron beam allowing capillary action of molten metals into the CF junction for high CF surface contact area. When the carbon fibers were electroplated with Ni prior to the partial welding, tensile stress, σ_b of Ni-plated carbon fiber [18-8/Ni–CF/Al] joint was 36 MPa, about 2.5 times larger than that without Ni plating [18-8/CF/Al] joint at 14 MPa. Fracture energy estimated by integrated area under the stress-strain curve was substantially improved in [18-8/Ni–CF/Al] over [18-8/CF/Al] joints. This shows the partial welding performed by spot electron beam allows joining metals with different melting temperatures with the carbon fiber junction. XRD (X-ray diffraction) tests revealed improvements in [18-8/Ni–CF/Al] joint are by Ni plating the carbon fibers minimizing typical brittle Al₄C₃ carbide formation by acting as a barrier, and rapid solution hardening by nickel addition. EPMA (Electron probe microanalysis) showed the Ni coating also acts to protect carbon fibers against impingement from the hot molten metals during welding, along with prevention of mutual diffusion between Al or Fe with carbon fibers increasing strength of the [18-8/Ni–CF/Al] joint.

Tensile stress-strain curves of joints of (18-8/Ni–CF/Al) with Ni-coating to carbon fibers, and (18-8/CF/Al) without Ni-coating, respectively, along with SEM image of Ni-coated carbon fibers. Results showed significant increase in tensile properties by Ni-coating.

1. Introduction

Increasingly strong and reliable joint technologies between different metals have always been required for advanced process technologies related to aerospace, high speed mover machines,¹⁾ along with ships, trains, building construction and other structures. One of the typical aerospace lightweight materials is aluminum (Al) which has electrical and thermal conductivity. Moreover, 18-8 stainless steel (∼18% chromium, ∼8% nickel) exhibits high strength, high resistance to oxidization at high temperature, and high electrical and thermal resistivity.²⁾ Both Al and 18-8 are widely used and exhibit high corrosion resistance at room temperature along with renewability and sustainability by just simple re-melting. Therefore, to make use of these properties an easy method to make a strong joint of the 18-8 and Al for practical applications is highly sought after.

Common joining methods used such as mechanical connection with bolts and rivets are relatively simple, however, disadvantages include stress intensity near the cutting hole, decay of fatigue strength, and increase in mass. Moreover, in welding brittle carbide or other chemical compound formation often occurs creating weak bonding.³⁾

Recently, in order to prevent mass gain by adding metal rivets or fasteners while controlling brittle compound formation by melting or welding, solid joint methods with metallic bonds have been reported. Three notable types include: friction-bonded joint;⁴⁾ plastic flow joint at Al-softening temperature;⁴⁾ and Al clad coils.^5,6) For friction-bonded joints, a maximum strength of 306 MPa is reported to have been obtained for friction-bonded joint of low carbon steel with about 0.10 mass%C to Al–Mg alloy (AA5083) at a friction time of 2 s under friction and forge pressures of 40 and 230 MPa, respectively. In addition, the plastic flow joint by forming projection of stainless steel to Al alloy at Al-softening temperature is reported to have been successfully developed and available for application to production lines utilized as sensors for automotive parts.⁷⁾ Joining Al clad coils to Ti or stainless steels, utilized for inductive heating (IH) cooking heater, have been successfully developed by semi-hot roll connecting method developed by Nippon Steel.^5,6)

On the other hand, for 18-8 steels and Al, the joint strength values by popular methods of welding and brazing at the joint interface are not sufficient due to formation of brittle intermetallic compounds.^3,8) For example, in Al-steel friction welded components, excessively thick intermetallic phase adversely affects bonding.³⁾ In addition, in Fe–Al solid solution, disordered Fe–Al phase represents the weakest points in a tensile test.⁸⁾ Furthermore, using high radiative heat it is difficult to weld the highly conductive 18-8 steel to highly thermally resistive Al. Success with radiation heat transfer welding has been mostly limited to plastic joints.^9–11)

Experimental results of this study showed σ-ε curves for an [18-8/Al] joint without CF junction could not be obtained by conventional welding. Therefore, with the high tensile strength (2 to 7 GPa) and tensile modulus, (200 to 500 GPa) of CF¹²⁾ we employ the CF junction for a [18-8/CF/Al] joint. The CF junction is a cross-weave cloth to easily set into the slits of the metals. Although the CF junction can generate stress concentrators by the spaces in the weave, the weave pattern is beneficial making composites less flaw sensitive giving significant improvement in mechanical properties of composites.¹³⁾ However, untreated CF has chemically inert surface with lack of bonding functional groups, poor interfacial wettability, and is not chemically stable with metals such as iron, thus limiting mechanical properties.¹⁴⁾

To remedy this, current research has focused on various methods of coating metals such as Cu, Ag, Co, Cr, Fe, Ti and Ni^15–18) onto CFs using physical vapor deposition (PVD), chemical vapor deposition (CVD), sol-gel, electroless plating, and electroplating.^19–23) A method to deposit Ni on CFs includes wire-mesh catalysis, which allows uniform and compact coating, with quality strongly dependent on the catalysis process.²⁴⁾ Carbonyl metal organic chemical vapor deposition (MOCVD) using Ni(CO) as a precursor has been used to deposit pure nickel onto CFs with good adhesion, resulting in increase in fracture strength of 34.9% and increased oxidation resistance over that of uncoated CFs.²⁵⁾

Methods of electroplating have been used to coat CFs with Ni.^12,26) In carbon fiber metal matrix composites (CF-MMC), coating CFs with ∼0.5 µm Ni thickness by an electroless plating method resulted in uniform distribution and Ni well bonded to CFs: tensile fracture strain was increased, while tensile stress was decreased in the Ni-coated fibers.¹²⁾ Continuous electroplating method has resulted in well adhered Ni to CFs and improvements in conductivity and oxidation resistance.²⁶⁾

Moreover, the literature reports^27–34) development of rapid melting prior to rapid solidification processes by: electron beam,^27–29) laser,³⁰⁾ direct current,^31,32) alternating current,³³⁾ and inductive magnetic field.³⁴⁾ In particular, the spot beam has advantages of highly localized beam, precisely controlled energy allowing rapid melting prior to rapid solidification, and vacuum atmosphere protecting molten metals from oxides and nitrides in the air.

The spot beam joining method has been applied to joining other materials: [Metal/Metal] and [Metal/Polymer] with a CF junction coated by Ni metal^28,29,35) taking advantage of the rapid melting-rapid solidification. For CF junctions, Ni is used as a coating for CFs to prevent: encroachment by molten metal into the carbon fibers at the high temperatures of welding; and brittle carbide formation at the CF/Metal interface from reaction between carbon and metal.^28,29,35) The Ni coating has the advantage of acting as a buffer preventing the CF encroachment during welding and brittle carbide formation, while promoting mutual diffusion between the Ni and metal matrix as a gradient absorbing energy during tensile testing increasing strength. For example, in a [Ti/Ni–CF/EpoxyCFRP] joint, Ni sputtering CF by a DC-magnetron sputtering device in low pressure argon atmosphere, prior to spot welding to Ti is reported to increase the impact value two times higher than that of either adhesive [Ti/Glue/Epoxy] joint with glue, or [Ti/Epoxy] joint without glue.³⁵⁾ The Ni coating to CF prevented embrittling TiC formation at the CF/Ti interface, along with enhancing the often imperfect contact between CF and Ti.³⁵⁾

The spot beam has been reported to be used with a “partial” welding method successfully developed for joining [Al/Ni–CF/Cu] and [Al/Ni–CF/Ti] joints.^27,28) This partial welding method has the advantage of melting the connecting metals separately around the CF junction especially when the metals have very different melting temperatures; while preventing brittle compound formation. A distinct advantage is creating a joint without use of drilling holes, fasteners, or adhesive. Since the CF junction is inserted into the “18-8 stainless steel half length” and adhered by rapid melting prior to rapid solidification, followed by that of “Al half length”, it is referred to here as “partial” welding method. Rapid solidification is used to prevent crystallization as is the case with amorphous metals by “freezing” a random structure before it can crystallize.³⁶⁾ Uniformity in noncrystalline structures eliminates grain boundary structure as in polycrystalline metals resulting in very high strengths and superior corrosion resistance.³⁶⁾ Rapidly solidified alloys may not be completely noncrystalline, but they are characteristically fine-grained, ∼0.5 µm compared to 50 µm for traditional alloys,³⁶⁾ the smaller grains lead to higher strengths.³⁶⁾ Melting temperatures reported for stainless steel and aluminum are very different, reported to be 1400° to 1450°C,³⁷⁾ and 660°C,³⁶⁾ respectively, hence the partial welding. For strong joints, the CF junctions afford an extremely large friction force by broad interface (CF: 6 µm-diameter), along with the rapid solidification to increase strength of the metals encasing the CF junction.

Up to now, spot welding used for partial welding process has not been reported for joining 18-8 and Al by Ni-coated CF junction. Therefore, with the goal of joining large articles for aerospace technology, this research demonstrates the construction of a new [18-8/Ni–CF/Al] joint from the difficult to join 18-8 stainless and Al via a Ni-coated CF junction by means of partial welding method. This is with the aim of creating a potential strong and light weight option for joining airplane structures with concern for the environment and safety. When applying the Ni coated CF junction to joining materials in a practical setting, carefulness is highly recommended to optimize for maximum safety and strength.

2. Experimental Procedure

2.1 Preparation of Ni-electroplated carbon fiber junction

Figure 1 shows SEM image of cross section of carbon fiber junction wrapped in Ni layer by electroplating referred to here as “Ni–CF”. As shown in the photograph, the width of Ni layer is homogeneously coated. Anode and cathode were Ni plate and carbon fibers, respectively. The carbon fibers (TORECA-M30SC, TORE; Tokyo, 6 µm) were coated by Ni electroplating (1.5 A, 7.0 V, 30 min) in 400 mL water solution containing 12 g boric acid, 100 g nickel sulfate, and 18 g nickel chloride at room temperature of 298 K.

Fig. 1

SEM-image of carbon fibers wrapped by Ni layer electroplated (Ni–CF).

2.2 Preparation of 18-8/Al joint connected by carbon fiber cloth junction

The 18-8 austenite stainless steel (18-8: SUS304 (Japan Industrial Standard) Niraco Co., Ltd. Tokyo) and high purity aluminum (Al: 99.3%-purity, Niraco Co., Ltd. Tokyo), respectively, are assembled separately prior to melting and welding on the graphite plate by using the rapid melting apparatus.²⁹⁾ The separate metal half length samples are remelted by movable electron beam spot (10 kV-constant electrical potential, 15 ± 5 mA-electrical current, 15 mm heat-effective spot diameter) controlled by magnetic field set in vacuum chamber (6.0 × 10⁻⁵ Pa-limited vacuum pressure, VA-8408, World Engineering Co., Ltd. Tokyo).

Figure 2 shows a schematic for [18-8/CF/Al] and [18-8/Ni–CF/Al] joint samples whose dimensions are: 60 × 10 × 3 mm. The 18-8 stainless steel and Al half lengths are shown (each 30 mm in length) totaling 60 mm for the entire joint sample. The inserted carbon fiber which acts as the junction (CF junction) is a cross-weave cloth with dimensions 20 × 10 × 0.3 mm, with 10 mm each into the 18-8 and Al. Here, “[18-8/CF/Al] joint” and “[18-8/Ni–CF/Al] joint” refer to the joints without, and with Ni coating to the CFs, respectively. Figure 3(a), (b) shows illustrations and photos of (a) [18-8/CF/Al] joint without Ni coating to CFs; and (b) [18-8/Ni–CF/Al] joint with Ni-coating, while Fig. 3(c) shows the two basic steps of the partial welding method employed. Step 1 is assembly of 18-8/CF part; while Step 2 is assembly of the projected CF of 18-8/CF part to Al. They are described in detail here:

Fig. 2

Schematic for [18-8/CF/Al] and [18-8/Ni–CF/Al] joint samples showing dimensions.

Fig. 3

Illustrations and photos of: (a) [18-8/CF/Al] and (b) [18-8/Ni–CF/Al] joint samples along with (c) assembly steps.

Step 1: Assembly of 18-8/CF part: To fabricate 18-8/CF part (bottom left, Fig. 3(c)), the CF or nickel plated CF cross-weave cloth (20 × 10 × 0.3 mm) is set into a cut slit (1.0 × 10 mm) into 18-8 stainless steel rod half length. The welding method is employed where half the length of the CF junction (10 mm) was contacted and wrapped with molten 18-8 by capillary phenomenon before solidification. The rapid melting by electron beam irradiation is performed at: V = 10 kV potential, I = 20 mA current for duration t = 30 s under P = 9.3 × 10⁻⁴ Pa residual gas pressure. The power term, t_p³⁸⁾ used is 10 s for the 18-8 stainless steel. Both size and motion of beam spot are controlled by magnetic field. Melting temperature range of 18-8 stainless steels is approximately 1400°C–1450°C.³⁷⁾

Step 2: Assembly of the projected CF of 18-8/CF part to Al: To join Al with the 18-8/CF part, the projected CF of 18-8/CF is set into a cut slit (1.0 × 10 mm) of Al rod. The partial welding method, of two metals melted around a CF junction separately is employed where the exposed length of CF cloth (10 mm) is dipped in molten pure Al then and solidified resulting in the finished [18-8/CF/Al] and [18-8/Ni–CF/Al] joint samples. Rapid melting parameters of electron beam irradiation are same as for the 18-8 except, I and duration t for the Al are set at: 10 mA and 60 s, respectively.

In both cases, the molten metal quickly penetrates to inter-spaces between carbon fibers. After that, rapid solidification due to thermal heat sink and low heat capacity occurs when the power is switched off. EB-melting prior to rapid solidification is performed under the same conditions for both half lengths of the samples. Melting temperature of pure aluminum is 660°C.³⁹⁾

2.3 Tensile test and characterization

When joint length, width and thickness are 10, 10 and 3 mm, tensile tests were carried out at a deformation rate of 1.0 mm/min at room temperature with an Autograph tensile tester (Shimadzu Model AG-10TE). Here, deformation rate refers to crosshead speed conforming with a previous study of Ni-coated CF reinforced Al connected with Ti rod, i.e. (Ti/NiCF/Al) joint.²⁸⁾

Stress-strain curves of the joint samples were obtained by using crosshead displacement and confirmed by using video recording device. Since the 18-8 and Al rods preferably deformed during tensile test, true stress–true strain curves were not adaptable because of heterogeneous deformation. The joint strength, σ_b (MPa) was obtained by the nominal stress-strain (σ-ε) curves.

For elemental and crystallographic analysis, sample cross sections were cut using a diamond blade to obtain flat surface. As shown in Fig. 4, the cut was taken 25 mm from the tail end of the length into the Al and 18-8 sections. Element mapping of C, Fe, Al, Ni and Cr was performed on cross section at carbon fiber and Al and 18-8 using an Electron probe micro-analyzer (EPMA-1610, 15 kV, 10 nA/Shimadzu, Kyoto).⁶⁾ X-ray diffraction (XRD) (Cu-Kα, MiniflexII, Rigaku, Tokyo)⁶⁾ was carried out using 10⁻³ deg/s scanning rate. Lattice structures of compounds were determined by the diffraction peaks evaluated by the ICDD (International Centre for Diffraction Data).

Fig. 4

Illustrations and photos of fractured: (a) [18-8/CF/Al], and (b) [18-8/Ni–CF/Al] joint samples. Dotted lines (25 mm from specimen ends) are cut locations for XRD and EPMA analyses.

3. Results

3.1 Tensile stress-strain curves of [18-8/Ni–CF/Al] and [18-8/CF/Al] joints

Figure 5 illustrates experimental tensile stress-strain (σ-ε) curves of the [18-8/Ni–CF/Al] and [18-8/CF/Al] joints, respectively. Although σ-ε curves for an [18-8/Al] joint without CF junction could not be experimentally obtained by conventional welding, implementation of CF junction assisted in generating maximum tensile strength (σ_b), its strain (ε_b), and fracture energy (U_T) estimated by the integrated area under the stress-strain curve is calculated by:   

\begin{equation} U_{\text{T}} = \int_{0}^{\varepsilon_{\text{f}}} \sigma \mathrm{d}\varepsilon \end{equation} (1)

thus confirming the 18-8 and Al could be joined together with CFs. The filled and unfilled data points in Fig. 5 show the σ-ε curves of [18-8/CF/Al] and [18-8/CF/Al] joints exhibiting σ_b of 14 and 36 MPa, respectively.

Fig. 5

Stress-strain curves of [18-8/Ni–CF/Al] joint with Ni coating to CFs and [18-8/CF/Al] joint without Ni coating, respectively.

The main result is electroplating the CFs with Ni-coating in the [18-8/Ni–CF/Al] joint prior to the partial welding improves tensile strength (σ_b), as well as its strain (ε_b), and fracture energy estimated by integration of σ-ε curves over that of [18-8/CF/Al] joint. Figure 5 shows the σ_b was improved nearly 2.5 times higher from 14 to 36 MPa and the ε_b was raised ∼63% from 0.016 to 0.026.

4. Discussion

4.1 Sample observation of fracture mechanisms

18-8 stainless steel and pure aluminum do not join well by welding due to brittle phase formation.^3,8) Tensile strength of composite joints is generally limited by the weakest links in the system in unison undergoing tension. The weakest sites are the 18-8/Al interface where brittle fracture occurs. Therefore, with the high tensile strength (2 to 7 GPa) and tensile modulus, (200 to 500 GPa) of CF¹²⁾ this research implements the strong CF junction to take on the majority of the load.

For observation of fracture mechanisms, Fig. 4 shows schematic illustrations and photos of fractured [18-8/CF/Al] (a) and [18-8/Ni–CF/Al] (b) joint samples, respectively (compare with Fig. 3). As expected, fracture occurs in ductile manner in the Al half length in both the [18-8/CF/Al] and [18-8/Ni–CF/Al] joints leaving the 18-8 stainless steel half length undamaged. This is due to higher tensile modulus (E) and ductility (% elongation at failure) of typical 18-8 stainless steel over alumninum.³⁶⁾ For example, typical 18-8 (304 stainless steel; E = 193 GPa and 40%) is stronger than typical aluminum (3003-H14 aluminum; 70 GPa and 8 to 16%).³⁶⁾

Experimental results show coating the CFs with Ni plays the dominant role in strengthening the joints. On the one hand, Fig. 4 shows the [18-8/CF/Al] joint exhibits a relatively straight, clean fracture of Al and CF cloth accompanied by slight CF pull-out along with ductile fracture of the Al (a). This is probably due to the CF surface damage in the form of impingement by the heat of the 18-8 and Al welding. On the other hand, the Ni-coated CF joint [18-8/Ni–CF/Al] joint exhibits significantly higher fracture surface area as shear fracture along the length of and within the CF cloth perpendicular to the cross section, in the same plane as the sample length and width (b). There is fiber fracture with gradual isolation of each carbon fiber from the cloth bundle observed on the ductile fracture surface of [18-8/Ni–CF/Al] joint Al half length. Therefore, fracture energy is increased in the [18-8/Ni–CF/Al] joint. Coating the CFs with Ni allows protection of the carbon fibers from the impingement, along with a strong Ni–Al interface for action of the CFs to take on more of the load over a much wider surface area raising the σ_b by nearly 3 times from 13 to 37 MPa.

4.2 Al-half length: Morphological discussion from XRD and EPMA data

Figure 6(a), (b) shows XRD data with ICDD standard to assess the state of Al–CF interface in the Al half length. The IDDC standards are shown above the XRD data. The Al cross section locations are indicated by dotted lines in Fig. 4(a).

Fig. 6

Al half length: XRD pattern of cross section of: (a) [18-8/CF/Al], and (b) [18-8/Ni–CF/Al] joints, together with ICDD standard.

Figure 6(a) indicates in the [18-8/CF/Al] joint, four XRD-peaks of brittle Al₄C₃ compound at diffraction angles 2θ = 56, 70, 77 and 90 deg are detected contributing to the lower tensile strength of the [18-8/CF/Al] joint. The XRD also detected pure Al (56 and 70 deg) and pure graphite (C) (28 and 47 deg) at Al–CF interface. Note the Al₄C₃ peaks are not accompanied by Al peaks at 77 and 90 deg, or are at much higher intensity than Al peaks at 56 and 70 deg. This clearly indicates formation of brittle Al₄C₃ in the [18-8/CF/Al] joint.

On the other hand, in Fig. 6(b) for the [18-8/Ni–CF/Al] joint, XRD-peaks show in addition to pure graphite at 2θ = 45, 65 and 78 deg, the CF protecting Ni–Al intermetallic compounds from the Ni electroplating are detected including ductile Al₃Ni₂ (45 deg), Al₃Ni (45 deg), AlNi₃ (45 deg), AlNi (45 and 65 deg). These act like a protective sheath or buffer for the CFs. The Al peak has higher intensity than accompanying low intensity Al₄C₃ peaks at 45, 65 and 78 deg although at 38 deg both Al and Al₄C₃ peaks are high intensity. These show the Ni–Al intermetallic compounds are induced by the mutual diffusion between the Ni and Al elements at Al/CF interface which blocks or hinders brittle Al₄C₃ carbide formation.

Figure 7(a), (b) shows EPMA mapping analysis of the Al half length cross section of the [18-8/CF/Al] (a) and [18-8/Ni–CF/Al] (b) joints, respectively. Figure 7(a) shows impinged carbon fibers with their relatively small diameters compared to those in Fig. 7(b). Figure 7(a) shows for [18-8/CF/Al] joint, smaller CF cross sectional areas accompanied by a diffusion gradient of C deeper into the Al is shown by increased green and yellow shades indicating formation of brittle Al₄C₃ compound. Mutually, it appears traces of Al are diffused into the CFs by increased light blue shade within CFs. On the other hand, Fig. 7(b) shows for [18-8/Ni–CF/Al] joint there is much less C diffusing into the Al due to the Ni coating acting as a diffusion barrier resulting in shorter C gradients into the Al indicated by increased blue areas around the CFs. This is the result of the Ni-coating forming Ni–Al chemical compounds around the CFs acting as protection barrier or sheath preventing mutual diffusion between the molten Al and CFs (see Fig. 6(b)). In addition, the Ni alloy coating prevents CF–Al debonding by mismatch in stiffness and by coefficient of thermal expansion difference.

Fig. 7

Al half length: EPMA mapping analysis of cross section of: (a) [18-8/CF/Al], and (b) [18-8/Ni–CF/Al] joints.

In this way, the [18-8/Ni–CF/Al] joint undergoes cloth shear fracture parallel to tensile direction; dissipating energy over a much wider surface area than that of the [18-8/CF/Al] joint, which exhibits a relatively straight, clean fracture perpendicular to the tensile direction with its considerably lower cross sectional area perpendicular to the test direction. The XRD and EPMA data show the Ni coating prevents molten Al impinging the CFs, while itself diffusing into the neighboring Al creating the strong bond increasing tensile strength and energy required for fracture.

4.3 18-8 stainless steel half length: Morphological discussion from XRD and EPMA data

Within the 18-8 half length on the other hand, Figs. 8(a), (b) and 9(a), (b) show XRD with ICDD standard, and EPMA mapping analysis of cross section at dotted lines in Fig. 4 for [18-8/CF/Al] (a) and [18-8/Ni–CF/Al] (b) joints, respectively.

Fig. 8

18-8 half length: XRD pattern of cross section of: (a) [18-8/CF/Al], and (b) [18-8/Ni–CF/Al] joints, together with ICDD standard.

Fig. 9

18-8 half length: EPMA mapping analysis of cross section of: (a) [18-8/CF/Al], and (b) [18-8/Ni–CF/Al] joints.

Since the high temperature of welding steel (about 2000 K) is much higher than that (about 1000 K) of Al, the impinging rate into CFs from steel would be extremely higher than that of Al. However, 18-8 steel having much higher mechanical strength than Al³⁶⁾ leads to the damage and fracture occurring in the Al half length rendering the 18-8 half length comparatively intact (Fig. 4).

As expected, XRD results in Fig. 8(b) show in the [18-8/Ni–CF/Al] joint 18-8 half length, Fe–Ni compounds are detected as sharp peaks at 2θ = 38 and 45 deg, a slight peak at 74 deg, and medium peak at 90 deg. In Fig. 8(a) in the [18-8/CF/Al] joint 18-8 half length, signals of α-Fe, γ-Fe are present at 2θ range of 20 to 35 deg. A sharp peak at 43 deg indicates Fe–Al compounds and Fe-carbides FeAl, Fe₂Al₅, Fe₃Al, FeC, Fe₂C, Fe₂C, Fe₃C, Cr₂C and C. Cr-carbide is detected as a medium intensity peak at 75 degrees.

As for morphology, specifically EPMA mapping results of the [18-8/CF/Al] joint in Fig. 9(a) show primary crystal grains 10–30 µm in size of α-Fe nucleate and grow between liquidus and solidus into residual molten steel. The γ-Fe grains generate at solidus. Both concentrations of carbon and Cr elements in γ-Fe are higher than those in α-Fe, whereas Ni concentration in γ-Fe is lower than that of α-Fe.

On the other hand, EPMA in Fig. 9(b) shows Ni-coating to CF of [18-8/Ni–CF/Al] joint generates the ductile Fe–Ni and Ni phases protecting the CFs, in addition to α-Fe, γ-Fe, Fe-carbides, Cr-carbide, Fe–Al compounds and graphite, as shown in XRD peaks of Fig. 8(b).

In addition, as shown in Fig. 9(b) EPMA mapping, carbon rich and poor phases are remarkably separated. The volume fraction of Al rich phase is larger than that of carbon rich phase. Cr-rich spot precipitation is observed in the Al rich phase.

Although the diameter of carbon fiber becomes smaller than that before making composites due to CF impingement with diffusion from CFs to steel matrix, damage and fracture is not observed in the 18-8 stainless steel half length in either the [18-8/CF/Al] or [18-8/Ni–CF/Al] joints due to its higher strength than Al.

To predict strengthening mechanism as a function of deformation region in the stress-strain curves of the joints, a “strain hardening” model is constructed in Fig. 10. Figure 5 shows the strain at ultimate tensile strength σ_b (ε_b) of [18-8/Ni–CF/Al] joint is much higher than that of [18-8/CF/Al] joint. Thus, constructing the extrapolation for the strain-hardening exponent of the [18-8/Al] joint should be estimated to predict the σ at each ε of the [18-8/CF/Al] and [18-8/Ni–CF/Al] joints. In general, the σ-ε curves of materials are expressed by the following equations:   

\begin{equation} \sigma = k\varepsilon^{\text{n}} \end{equation} (2)

  

\begin{equation} \log \sigma = \log k + n \log \varepsilon \end{equation} (3)

Here n is strain hardening index and k (MPa) is the strength coefficient. The strain hardening index is n = 0 for a perfectly plastic solid, and n = 1 for a perfectly elastic solid.

Fig. 10

log₁₀ σ-log₁₀ ε curves of [18-8/CF/Al] and [18-8/Ni–CF/Al] joints calculated from Fig. 5 showing strain hardening index, n indicating deformation mechanism regions.

Figure 10 shows the base-10 logarithmic stress (σ)–strain (ε) curves (log₁₀ σ-log₁₀ ε) of [18-8/Ni–CF/Al] and [18-8/CF/Al] joints, respectively from data in Fig. 5. The plots are separated into 2 main regions of fracture mode: elastic and plastic. Plastic region is separated into sub-regions 2-i where i = 1 to 3. Resulting n values are depicted in Fig. 10 and summarized in Table 1.

Table 1 Deformation regions with strain hardening coefficients, n for (18-8/CF/Al) and (18-8/Ni–CF/Al) joints.

Region 1: The homogeneous deformation of Region 1 (n = ∼1) obeying Hooke’s Law is mainly caused by reversible elastic deformation of series connection of Al rod, CF junction, and 18-8 rod. In Region 1, CFs take on most of the load transferring by adhesion to the surrounding metals Al and 18-8.

Without Ni coating to CFs, the [18-8/CF/Al] joint soon transitions into Region 2 with plastic deformation due to low elastic strain limit of the brittle Al–C compounds formed at the fiber-matrix interface, weakened CFs by impingement, and ductile deformation of the Al half length. However, the Ni coating to CFs in [18-8/Ni–CF/Al] joint allows elastic behavior of Region 1 up to much higher strains.

Region 2-1: Here, n is reduced by transition from elastic to plastic deformation. It shows Ni plating to the CFs significantly improves the onset stress and strain of plastic deformation in the [18-8/Ni–CF/Al] over the [18-8/CF/Al] joint. Both joints apparently have the same strain hardening index of n = ∼1/2 from similar deformation mechanisms. They are categorized here into 4 modes:

Mode 1) In both [18-8/CF/Al] and [18-8/Ni–CF/Al] joints, irreversible damage or debonding at carbon fiber-Al matrix interface occurs from the onset of Region 2-1. However, in [18-8/CF/Al] joint Region 2-1 initiates at lower strains due to low elastic strain limit of the brittle Al–C compounds present. On the other hand, in [18-8/Ni–CF/Al] joint Region 2-1 occurs at high strains above the elastic strain limit of the CF–Ni-coating-Al interface. The Ni prevents brittle Al–C compound formation for a stronger cohesion between CFs and metal.

Mode 2) In both joints carbon fiber breakage occurs in Region 2-1, however in [18-8/CF/Al] joint CF breakage occurs at low strains by weakened bare CFs by impingement, while in [18-8/Ni–CF/Al] joint CF breakage is prevented up to high strains by the Ni coating to the CFs resulting in Al–Ni compound formation acting as a buffer between the CF and metal.

Modes 3, 4) In both joints, damage at 18-8/Al interface layer (Mode 3) and ductile deformation of the Al half length (Mode 4) occurs simultaneously.

When the plastic deformation also irreversibly enlarges the resistance to stress of aluminum, considerable plastic deformation with relaxation with work-hardening²⁸⁾ occurs. The work-hardening and simultaneous relaxation are probably caused by cross-slip in Al crystalline generated by tensile loading.

Region 2-2: The [18-8/CF/Al] joint without Ni coating exhibits strain hardening coefficient n increasing to about 2 designated here as Region 2-2. After the early ductile deformation of Al in Region 2-1, the cloth CFs apparently take on more of the load being stretched raising the strain hardening coefficient of the system significantly to n > 2, above the initial n = ∼1.0 in Region 1. Another factor probably raising the n is work-hardening of Al matrix containing the partially fractured impinged CFs by Al texture forming preferred orientation parallel to the tensile direction as it undergoes the tensile deformation.

In contrast, the Ni containing [18-8/Ni–CF/Al] joint exhibits a much more stable deformation and does not undergo Region 2-2 due to action of the Ni enhancing adhesion of CFs and Al. Although the [18-8/Ni–CF/Al] joint does not exhibit the high n > 2, it undergoes elastic deformation (Region 1) to a much higher load than the [18-8/CF/Al] joint.

Region 2-3: This is a short strain range characterizing the final damage just prior to fracture with slight work-hardening (n < 0.1). Damage is in the form of heterogeneous plastic deformation of necking and ductile fracture in the Al, occurring simultaneously with CF pullout and breakage. The [18-8/CF/Al] joint exhibits a relatively straight, clean fracture perpendicular to the tensile direction (Fig. 4(a)) with n = 0.0419. On the other hand, Fig. 4(b) shows the [18-8/Ni–CF/Al] joint exhibits the strengthening mechanism of peeling across a much wider surface area parallel to the tensile direction exhibiting extremely higher fracture energy than the [18-8/CF/Al] joint. Note Fig. 10 shows Region 2-3 strain range for [18-8/Ni–CF/Al] joint is slightly longer than that for [18-8/CF/Al] joint; and the [18-8/Ni–CF/Al] joint has higher n = 0.0957 than that of [18-8/CF/Al] joint at n = 0.0419.

Region 2-3 ends with the fracture point at ultimate tensile stress, σ_b, in both joints.

Although joint samples connecting 18-8 with Al were made successfully using CFs, Ni-coating to CF controls the formation of brittle Al carbides and attains significantly higher strain and stress in [18-8/Ni–CF/Al] joint. Coating the CFs with Ni increases joining strength by strain hardening well beyond the strength of [18-8/CF/Al] joint. A strength hierarchy can be constructed for the increased strain hardening: [18-8/Al] joint → [18-8/CF/Al] joint CF reinforcement → [18-8/Ni–CF/Al] joint Ni-coated CF reinforcement.

The “strain hardening” model in Fig. 10 and Table 1 aims to be a useful tool to predict deformation mechanisms in CF reinforced [18-8/Al] joints or similar [Metal/CF/Metal] joints. Carefulness for maximum safety is advised when constructing [18-8/Ni–CF/Al] joints for aircraft applications.

5. Conclusions

• (1)    A novel [18-8/Ni–CF/Al] joint of 18-8 stainless steel and aluminum connected by nickel-coated carbon fiber cross-weave junction, utilizing extremely large friction force by broad interface of 6 µm diameter CF has been successfully developed for potential aerospace application. This was done by a partial welding process of the CF junction to 18-8 stainless steel and Al separately taking advantage of the different welding temperatures for Al, and stainless steel. Both the 18-8 stainless steel, and Al were remelted by movable electron beam spot controlled by magnetic field set on a graphite plate in vacuum chamber. The partial welding has the advantage of melting the connecting different metals separately around the CF junction while the junction itself provides broad surface area from the 6 µm fibers. Electroplating the CFs with Ni prior to welding creating a [18-8/Ni–CF/Al] joint increasing the tensile properties significantly over [18-8/CF/Al] joint.
• (2)    While the [18-8/CF/Al] joint exhibited maximum tensile stress σ_b and strain ε_b of 14 MPa and 0.016, respectively, electroplating the CFs with Ni prior to the partial welding in [18-8/Ni–CF/Al] joint increased the σ_b nearly 2.5 times to 36 MPa and the ε_b ∼ 63% from 0.016 to 0.026. In addition, toughness as area under the stress-strain curve was increased in the [18-8/Ni–CF/Al] joint.
• (3)    Optical observation of the joints showed the [18-8/CF/Al] joint exhibited a relatively straight, clean fracture perpendicular to the tensile testing direction of Al and CF cloth accompanied by slight CF pull-out along with ductile fracture of the Al in the form of necking. On the other hand, the Ni-coated CF [18-8/Ni–CF/Al] joint in addition exhibited significantly higher fracture surface area as shear fracture along the length of and within the CF cloth perpendicular to the cross section, in the same plane as the sample length and width.
• (4)    XRD results of cross section in the Al half length showed in the [18-8/CF/Al] joint four XRD peaks of brittle Al₄C₃ compound were detected causing weaker CF–Al interface. However, in the [18-8/Ni–CF/Al] joint XRD detected Ni–Al intermetallic compounds from the Ni electroplating including ductile Al₃Ni₂, Al₃Ni, AlNi₃, and AlNi increasing tensile strength.
• (5)    Moreover, EPMA results in the Al half length showed the [18-8/Ni–CF/Al] joint had larger CF cross sectional areas than the [18-8/CF/Al] joint. This indicates the Ni-coating acted to protect the CFs in the [18-8/Ni–CF/Al] joint from impingement by the hot metal during welding. Also, the EPMA results in the Al half length show the Ni-coating prevents mutual diffusion of species which could be detrimental to strength. The [18-8/CF/Al] joint had a diffusion gradient of C deeper into the Al and mutually, traces of Al diffused into the CFs much more than in the [18-8/Ni–CF/Al] joint.
• (6)    Coating the CFs with Ni allows protection of the CFs from impingement by the high temperature of the spot beam, along with a strong Ni–Al interface for action of the CFs to take on more of the load over a much wider surface area raising the σ_b by nearly 3 times from 13 to 37 MPa.
• (7)    A “strain hardening” model is constructed aiming to be useful to predict deformation mechanisms as a function of deformation region in the stress-strain curves in the [18-8/CF/Al] and [18-8/Ni–CF/Al] joints.
• (8)    The Ni coating is highly recommended to act as protective sheath or buffer between the CFs and metals from brittle carbide formation and impingement from hot metal during welding for a stronger joint. The Ni coating contributes to the higher tensile properties of the [18-8/Ni–CF/Al] joint.

Acknowledgments

The authors wish to thank Mr. Yasuo Miyamoto of Tokai University for his excellent help with EPMA (Electron Probe Microanalysis). Sincere gratitude goes to the Matsumae International Foundation (MIF) and the Japan Society for the promotion of Science (JSPS) Core-to-Core Program for their great support.

REFERENCES

• 1)   S.  Katayama: J. Light Metal Welding & Construction 42 (2004) 16–25 (in Japanese).
• 2)   K.  Nonaka,  H.  Fujii and  H.  Nakajima: J. Japan Inst. Metals 64 (2000) 85–94 (in Japanese).
• 3)   M.  Yılmaz,  M.  Çöl and  M.  Acet: Mater. Charact. 49 (2002) 421–429.
• 4)   N.  Yamamoto,  M.  Takahashi,  K.  Ikeuchi and  M.  Aritoshi: Mater. Trans. 45 (2004) 296–299.
• 5)   K.  Hamada,  K.  Isaka,  D.  Doi,  Y.  Yonemitsu and  S.  Iwasaki: Tetsu-to-Hagané 88 (2002) 66–72 (in Japanese).
• 6)  T. Okui, Y. Yonemitsu and K. Yoshida: Technical Review of Nippon steel & Sumikin Engineering Co. LTD. No. 396 (UDC 669 . 71 - 419 . 4 : 669 . 14 . 018 . 8 : 669 . 295) (2013) 106–110 (in Japanese).
• 7)   H.  Shirai,  M.  Mochizuki and  M.  Toyota: J. Light Metal Welding 48 (2010) 286–292 (in Japanese).
• 8)   P.  Šesták,  M.  Friák,  D.  Holec,  M.  Všianská and  M.  Šob: Nanomaterials 8 (2018) 873–892.
• 9)   Y.  Kurosaki: J. Quant. Spectrosc. Radiat. Transf. 93 (2005) 25–41.
• 10)   R.K.  Nath,  P.  Maji and  J.D.  Barma: J. Braz. Soc. Mech. Sci. Eng. 41 (2019) 553.
• 11)   D.Y.  Kim and  K.Y.  Yi: Trans. Korean Inst. Electr. Eng. 63 (2014) 1098–1103.
• 12)   Y.  Xue,  W.  Chen,  Q.  Zhao and  Y.-Q.  Fu: Surf. Coat. Tech. 358 (2019) 409–415.
• 13)   B.  Muralidhara,  S.P.  Kumaresh Babu and  B.  Suresha: Materials Today: Proceedings 22 (2020) 1755–1764.
• 14)   N.R.  Calderon,  R.  Voytovych,  J.  Narciso and  N.  Eustathopoulos: J. Mater. Sci. 45 (2010) 2150–2156.
• 15)   Y.  Wan,  J.  Xiao,  C.  Li,  G.  Xiong,  R.  Guo,  L.  Li,  M.  Han and  H.  Luo: J. Magn. Magn. Mater. 399 (2016) 252–259.
• 16)   P.  Baumli,  J.  Sychev,  I.  Budai,  J.T.  Szabo and  G.  Kaptay: Compos. Part A 44 (2013) 47–50.
• 17)   X.W.  Zhao,  C.G.  Zang,  Q.K.  Ma,  Y.Q.  Wen and  Q.J.  Jiao: J. Mater. Sci. 51 (2016) 4088–4095.
• 18)   X.G.  Cao,  C.L.  Liang and  S.M.  Xia: Mater. Sci. Forum 847 (2016) 137–142.
• 19)   Z.  Guo,  L.  Sang,  Z.  Wang,  Q.  Chen,  L.  Yang and  Z.  Liu: Surf. Coat. Tech. 307 (2016) 1059–1064.
• 20)   J.A.  Garlow,  L.K.  Barrett,  L.  Wu,  K.  Kisslinger,  Y.  Zhu and  J.F.  Pulecio: Sci. Rep. 6 (2016) 19804.
• 21)   C.  Sun,  S.  Chen and  Z.  Li: Appl. Surf. Sci. Part A 427 (2018) 476–484.
• 22)   Y.  Feng,  G.E.  Mcguire,  O.A.  Shenderova,  H.  Ke and  S.L.  Burkett: Thin Solid Films 615 (2016) 116–121.
• 23)   B.  Abbasipour,  S.M.  Monirvaghefi and  B.  Niroumand: Met. Mater. Int. 18 (2012) 1015–1021.
• 24)   W.Q.  Zeng,  W.M.  Zeng,  Z.Z.  Gao and  Y.L.  Ma: Surf. Technol. 46 (2017) 276–281.
• 25)   Y.  Li,  J.H.  Nie,  N.  Li,  X.Q.  Liu and  C.C.  Jia: J. Funct. Mater. 43 (2012) 1688–1691.
• 26)   X.X.  Lv,  C.X.  Lv,  Y.  Yang,  F.  An,  Y.L.  Ji and  F.  He: New Chem. Mater. 39 (2011) 89–91.
• 27)   Y.  Nishi and  K.  Tanaka: Mater. Trans. 47 (2006) 1821–1825.
• 28)   N.  Harigae and  Y.  Nishi: J. Japan Inst. Metals 75 (2011) 235–242 (in Japanese).
• 29)   K.  Shiraishi,  S.  Inui,  S.  Ishii,  Y.  Matsumura and  Y.  Nishi: Mater. Trans. 55 (2014) 1564–1567.
• 30)   H.  Ozaki and  M.  Kutsuna: Quarterly J. Japan Welding Society 25 (2007) 473–479 (in Japanese).
• 31)   A.  Mizutani,  N.  Uchida,  K.  Oguri and  Y.  Nishi: J. Japan Inst. Metals 68 (2004) 158–161 (in Japanese).
• 32)   Y.  Harada,  Y.  Yamada and  S.  Kumai: J. Inst. Light Metals 66 (2016) 489–496 (in Japanese).
• 33)   H.J.  Park,  S.  Rhee,  M.J.  Kang and  D.C.  Kim: Mater. Trans. 50 (2009) 2314–2317.
• 34)   J.-Y.  Shim,  I.-S.  Kim,  M.-J.  Kang,  I.-J.  Kim,  K.-J.  Lee and  B.-Y.  Kang: Mater. Trans. 52 (2011) 999–1002.
• 35)   Y.  Nishi,  S.  Ishii,  S.  Inui,  A.  Kasai and  M.C.  Faudree: Mater. Trans. 55 (2014) 323–326.
• 36)  J.F. Shackelford: Introduction to Materials Science for Engineers, 5th ed., (Prentice Hall International, London, 2000) pp. 136, 137, 194, 412.
• 37)  Handbook of Stainless Steel, Outokumpu, (Outokumpu Oyj, Finland, 2013) p. 22.
• 38)   H.  Hasegawa,  M.C.  Faudree,  Y.  Enomoto,  S.  Takase,  H.  Kimura,  A.  Tonegawa,  Y.  Matsumura,  I.  Jimbo,  M.  Salvia and  Y.  Nishi: Mater. Trans. 58 (2017) 1606–1615.
• 39)  A. James and M. Lord: Macmillan’s Chemical and Physical Data, (The Macmillan Press, Ltd., London and Basingstoke, 1992) p. 285.