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Oxidation of stainless steel in vacuum and evolution of surface oxide scales during hot-compression bonding
Accepted Manuscript Title: Oxidation of stainless steel in vacuum and evolution of surface oxide scales during hot-compression bonding Authors: Bijun Xie, Mingyue Sun, Bin Xu, Chunyang Wang, Haiyang Jiang, Dianzhong Li, Yiyi Li PII: DOI: Reference:
S0010-938X(18)31547-6 https://doi.org/10.1016/j.corsci.2018.11.001 CS 7757
To appear in: Received date: Revised date: Accepted date:
17 September 2018 5 November 2018 10 November 2018
Please cite this article as: Xie B, Sun M, Xu B, Wang C, Jiang H, Li D, Li Y, Oxidation of stainless steel in vacuum and evolution of surface oxide scales during hot-compression bonding, Corrosion Science (2018), https://doi.org/10.1016/j.corsci.2018.11.001 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Oxidation of stainless steel in vacuum and evolution of surface oxide scales during hot-compression bonding Bijun Xiea,b, Mingyue Suna,d,*, Bin Xua,d, Chunyang Wanga,b, Haiyang Jiangc,
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Shenyang National Laboratory for Materials Science, Institute of Metal Research,
b
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Chinese Academy of Sciences, Shenyang 110016, China
c
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Dianzhong Lia, Yiyi Lia
University of Chinese Academy of Sciences, Beijing 100049, China
School of Metallurgy and Materials Engineering, Jiangsu University of Science and
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Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal
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d
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Technology, Zhenjiang 212003, China
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Research, Chinese Academy of Sciences, Shenyang 110016, China
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*Corresponding author. Tel.: +86 24 83978839; E-mail: [email protected]
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Highlights
The oxidation behaviors of austenite stainless steel in vacuum were investigated.
Thick oxide scale with outer MnFe2O4 and inner FeCr2O4 formed in a low vacuum.
Complex surface oxide scale transformed to MnCr2O4 after hot compression
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bonding.
Oxide scale with outer Fe3O4 and inner (Fe, Cr)-rich layer formed in a high vacuum.
Nanoscale-thick surface oxide scale decomposed during hot compression bonding.
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Abstract
The oxidation behaviour of austenite stainless steel in vacuum and the evolution of oxide scales during vacuum hot-compression bonding were investigated combining
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XPS and TEM. The results indicate that thick duplex oxide scale with outer MnFe2O4 and inner FeCr2O4 layer was formed in a low vacuum and the complex oxide scale
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transformed into interfacial oxides MnCr2O4 remaining at the bonding interface after
hot-compression bonding at 1200 °C. While in a high vacuum, nanoscale-thick oxide scale with outer Fe3O4 and inner Fe- and Cr-rich layer was observed and it decomposed
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during the subsequent hot-compression bonding procedure, leading to a well-bonded
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interface.
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Keywords: Stainless steel; Vacuum oxidation; XPS; TEM; Hot-compression bonding
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1. Introduction
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Since stainless steel clad plates simultaneously exhibit good mechanical properties and corrosion resistance, they have attracted increasing attention for various
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engineering applications, e.g. heat exchangers, nuclear power equipment, transferring pipes, vessels, and automobiles [1-3]. Stainless steel clad plates are layered structures
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that are generally fabricated by bonding stainless steel plates to other metal plates, such as carbon steel, low-alloy steel, titanium alloy, or aluminium alloy [4-7]. Different technologies are used to manufacture stainless steel clad plates, e.g. diffusion bonding, explosion welding, and roll bonding [4, 8-10]. Among all technologies, hot roll bonding is the most effective and suitable manufacturing procedure for clad plates during which 2
the metal plates are joined under high temperatures [11]. However, regarding the conventional hot roll bonding technique, as metal surfaces significantly react with atmospheric oxygen, oxide scales form on the contacting surfaces during heating. This deteriorates the bonding quality of the clad plates. Therefore, vacuum rolling cladding
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(VRC) is applied to obtain better bonding effect during which the contacting surfaces of the metal plates are cleaned by mechanical polishing and chemical etching to remove
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barriers, and then bonded in a vacuum. However the fact is that oxidation at the bonding
interface is still inevitable even under vacuum conditions. For example, Xie et al. [12]
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showed that the Mn–Si–O oxide particles were formed at the clad interface during VRC
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process in a vacuum of 10-1 Pa. Zhu et al. [13] reported the formation of interfacial
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oxide at 316L stainless steel cladding interface and indicated that the interfacial oxide
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is the limiting factor for interface bonding. Actually, these interfacial oxides at the clad
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interface are determined by the initial surface oxidation of the contacting surfaces during initial holding process under high temperatures and vacuum conditions before
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rolling cladding. Therefore, it’s essential to investigate the initial high-temperature
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oxidation of the contacting surfaces under vacuum conditions in order to control and eliminate interfacial oxides to acquire better bonding effect. This is particularly necessary for stainless steel clad plates with compact and stable oxide scales on the
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stainless steel surface. The oxidation behaviour of stainless steel in super critical water, water vapour, steam and ambient air at high temperature has been extensively studied [14-18]. In all these cases, it has been comprehensively reported that the oxide scale has a duplex structure, 3
composed of Fe and Cr oxides. The outer layer of the scale is predominantly composed of Fe-rich oxide, whereas the inner layer consists of Cr-rich oxide. Furthermore, Li et al. [19] reported the early oxidation mechanism of stainless steel TP347H in supercritical water. Matthews et al. [20] investigated the intergranular oxidation of 316L
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stainless steel in the PWR water environment and suggested that an epitaxial relationship existed between the metal and the inner Cr-rich oxide. Swaminathan et al.
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[21, 22] reported that the oxide scale formed in ambient air was mainly composed of
mixed oxides such as FeCr2O4 and MnCr2O4 along with binary oxides with Fe, Cr, and
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Mn. Col and Cheng et al. [23, 24] focused on the breakaway oxidation behaviour of
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stainless steel at high temperature. It is reported that the breakaway oxidation behaviour
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was significantly influenced by the microstructure and the composition of the oxide
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scale and the occurrence of breakaway oxidation was related to a local conversion of
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Cr2O3 in less protective FeCr2O4. Ostwald et al. [25] found that the oxide scale formed on stainless steel in H2–2.5%H2O atmosphere was protective and contained high Cr,
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Mn and Si contents, while in N2-20.5%O2 atmosphere, the oxide scale was Fe- and Ni-
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containing oxides. Nevertheless, to the best of our knowledge, systematic investigations on the high-temperature oxidation of stainless steel under vacuum conditions are scarce. The objective of this work is to investigate the oxidation behaviour of austenite
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stainless steel in vacuum and the evolution of oxide scales during vacuum hotcompression bonding. By using X-ray photoelectron spectroscopy (XPS) and transmission electron microscopy (TEM), we systematically studied the surface chemistry and structure of oxide scales formed on the surface of 316LN stainless steel 4
under two different vacuum conditions (10-1 Torr and 10-4 Torr). Moreover, the evolution of the surface oxide scales during hot-compression bonding in corresponding vacuum environments was investigated. The results show that the initial surface complex oxide scale transformed into an interfacial MnCr2O4 layer during hot-
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compression bonding in a low vacuum, whereas the nanoscale-thick surface oxide scale formed in a high vacuum has decomposed after hot-compression bonding. The
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systematic investigation of the surface oxidation and evolution of oxide scales during
vacuum hot-compression bonding provide new insight into high-temperature oxidation
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in vacuum and may have implications in manufacturing of high-quality bonding joints.
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2. Experimental method
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Austenite stainless steel (316LN) with a chemical composition (in wt.%) of 0.012 C,
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0.12 N, 0.34 Si, 0.003 S, 0.018 P, 12.83 Ni, 1.41 Mn, 2.18 Mo, 16.93 Cr, 0.016 Al, 0.04 Cu, 0.059 Nb and balance Fe was used as raw material. Square samples with the
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dimensions of 8*8*6 mm were cut from the raw material. The test surfaces were wet-
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ground and buff-polished with diamond paste. The buff-polished surfaces are named as “native surface”. Two groups of square samples were heated in the Gleeble-3800
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thermal simulator under different vacuum conditions, i.e. 10-4 Torr and 10-1 Torr, respectively. During the heat treatment, the samples were first heated to 1200 °C and maintained for 5 min. Afterwards, the samples were cooled to reach 30 °C and then removed from the vacuum environment. The surfaces of the treated samples are called “oxidised surface I” and “oxidised surface II” for two vacuum environments of 10-4 5
Torr and 10-1 Torr, respectively. After this treatment, the “native surface”, “oxidised surface I”, and “oxidised surface II” were investigated to understand the oxidation of the contacting surfaces under different vacuum conditions using XPS and TEM. Furthermore, the hot-compression bonding experiments (Fig. 1a) were performed in the
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Gleeble-3800 with vacuum degree of 10-4 Torr and 10-1 Torr at the strain rate of 0.1. The interfacial characteristics of two kinds of bonding interfaces were analysed by
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scanning electron microscope (SEM) and TEM.
The XPS measurements were performed using a PHI Quantera II equipped with a
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monochromatic Al Kα (1486.6 eV) X-ray source. To determine the chemistry of the
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oxide scales formed on the “native surface”, “oxidised surface I” and “oxidised surface
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II”, XPS survey scans and high-energy resolution spectra of characteristic peaks of
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elements such as Fe 2p, Cr 2p, Mn 2p, Si 2p, C 1s, and O 1s were acquired. In order to
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determine the elemental distribution in depth from the surface, i.e. the oxide chemistry in the sub-surface zones, the samples were sputtered with Ar+ ions with an energy of 2
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keV. The elemental spectra were fitted using Casa XPS software with Gaussian–
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Lorentzian line shapes and the Shirley-type background.
Fig. 1. (a) Schematic of hot-compression bonding process of two specimens under different vacuum conditions. (b) Preparation of transmission electron microscopy 6
(TEM) sample by focused ion beam (FIB) and (c) thinning of the sample lamella with the oxide scale on it using FIB.
The TEM sample containing the oxide scale was prepared and extracted using a
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focused ion beam (FIB) (FEI Helios NanoLab 650 DualBeam) with the in-situ loft-out (INLO) method, as shown in Fig. 1b. After which, the sample was further thinned to
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electron transparency by ion milling (Fig. 1c). Detailed structural and elemental
characterizations of the oxide scales and interfacial oxides were conducted using an FEI
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Tecnai G2 F20 transmission electron microscope attached with a high-angle annular
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dark-field detector and an EDS system with an acceleration voltage of 200 kV.
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3. Results
3.1 XPS analyses of oxide scales formed on “native surface”, “oxidised surface I”, and “oxidised surface II”
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3.1.1 Depth analyses
Fig. 2. Variations in chemical composition with respect to depth: (a) “native surface”,
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(b) “oxidised surface I” (treated in a high vacuum of 10-4 Torr), (c) “oxidised surface II” (treated in a low vacuum of 10-1 Torr).
The variations in the chemical composition with depth were obtained on three types of oxidised surfaces using XPS (Fig. 2). The XPS oxygen profile in Fig. 2a shows that 8
the oxide scale on the “native surface” is approximately 4 nm thick and there is a decreasing oxygen gradient, after which the composition of the matrix starts. Fe and Cr enrichment is visible at the surface (top 4 nm). As for the “oxidised surface I” (treated in a high vacuum of 10-4 Torr), the oxide scale thickness is approximately 30 nm and a
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decreasing oxygen gradient over ~50 nm is observed (Fig. 2b). It can be seen that the Fe and Cr enrichment is obvious throughout the oxide scale. Fig. 2c shows an overview
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composition of the top 3.6 μm of “oxidised surface II” (treated in a low vacuum of 10 Torr) and this oxide scale thickness is approximately 500 nm. Further, an internal
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oxidation zone (IOZ) of approximately 2.5 μm exists, after which the composition of
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the matrix starts. The initial Fe and Mn enrichment around 50 nm in depth in Fig. 2c
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indicates that Fe- and Mn-rich oxides are likely formed on the top surface. The obvious
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Cr enrichment at the surface (top 500 nm) demonstrates the presence of Cr-rich oxides.
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Besides, the Si enrichment is also observed in the top 500 nm.
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3.1.2 Surface chemistry
Fig. 3. XPS survey and high-energy resolution spectra obtained at different sputtering
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depths of “native surface”: (a) survey spectra, (b) Cr 2p spectra and (c) Fe 2p spectra.
Fig. 3 presents the XPS survey scan and high-resolution spectra of “native surface”.
The depth profiles were obtained for different sputtering-time intervals to assess the
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chemical state of elements at different depths of the oxide scales. The outmost surface spectra Cr 2p3/2 and Fe 2p3/2 (Figs. 3b and c) have binding energies of 576.1 eV and 710.37 eV, corresponding to Cr oxide and Fe oxide, respectively [26]. This confirms that binary oxides of Cr and Fe are formed on the initial “native surface”. Additionally, 10
tiny shoulders at 573.46 eV (Cr 2p3/2), 706.33 eV (Fe 2p3/2) and 638.8 eV (Mn 2p3/2) indicate their metallic state [26] and these metallic spectral peaks increase with increasing sputter depth. After sputtering down to 5 nm, only metallic spectral peaks appear (Figs. 3a-c), which indicates that the thickness of the oxide scale formed on the
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“native surface” is below 5 nm.
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Fig. 4. XPS survey and high-resolution spectra obtained at different sputtering depths
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of “oxidised surface I” (oxidised in a high vacuum of 10-4 Torr): (a) survey spectra, (b) Si 2p spectra, (c) Cr 2p spectra, (d) Fe 2p spectra, (e) Curve fitted Fe 2p3/2 spectra
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obtained at sputtering depth of 0 nm, (f) Deconvolution of Fe 2p3/2 spectra obtained at sputtering depth of 12 nm.
In-depth distributions of the XPS survey and high-resolution spectra of Si 2p, Cr 2p and Fe 2p of “oxidised surface I” are presented in Fig. 4. According to the survey 11
spectra (Fig. 4a), The Mn spectral peak at 638.6 eV (Mn 2p3/2) was only observed in the deeper region, indicating its metallic state. A Si 2p spectral peak was also detected, and the Si 2p3/2 spectral peak (103.4 eV) [26] only appeared at the surface regions, corresponding to the oxidized Si. No spectral peaks were observed at deeper regions
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owing to the low Si content in the matrix (Fig. 4b). Regarding depth-resolved Cr 2p spectra, the binding energy of 576.6 eV for Cr 2p3/2 spectrum in Fig. 4c corresponds to
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Cr oxide [26, 27]. Similarly, after sputtering down for a certain depth, a metallic Cr peak appeared at 573.8 eV (Cr 2p3/2) [27]. The depth-resolved Fe 2p3/2 spectra in Fig.
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4d positioned the binding energy at 710.5 eV (at the surface), indicating that Fe exists
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in oxidized state of Fe3O4 at the surface [26] (Fig. 4e). With increasing sputtering depth,
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Fe 2p3/2 peak shifted towards a smaller binding energy and broadened (at the sub-
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surface), suggesting a mixed Fe state. Fig. 4f illustrates the deconvolution of Fe 2p 3/2
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spectrum obtained after 12 nm sputtering into the oxide scale formed on “oxidised surface I”. Three components were necessitated to fit the Fe 2p3/2 transition. The higher
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binding energy component at 710.5 eV is attributed to Fe3O4 [26], and the lower binding
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energy component at 709 eV is due to Fe2+ which may be from FeO or (Fe,Cr)3O4 oxides [26, 28]. The binding energy component at 706.8 eV belongs to metallic Fe [26], and the intensity of the metallic spectral peak increases with increasing sputter depth
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(Fig. 4d).
Fig. 5 shows the XPS survey and high-resolution (Si 2p, Cr 2p, Fe 2p and Mn 2p) spectra obtained on the “oxidised surface II” at different depths. The in-depth distributions of Si 2p and Cr 2p spectra (Figs. 5b and c) have similar variations as that 12
of the “oxidised surface I” (Figs. 4b and c). As for Fe 2p spectra (Fig. 5d), the binding energy of Fe 2p3/2 spectra range from 710.7 eV (at the surface) to 709.4 eV (at the subsurface) corresponds to Fe oxides [26]. This oxide spectral peak gradually weakens with increasing sputter depth towards the matrix and the metallic Fe peak at 706.9 eV
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(Fe 2p3/2) appears (Fig. 5d) [26]. It is noteworthy that the intensity of Fe 2p3/2 peak is weak at the depth of 236 nm, which indicates that Fe oxide content is low at this region
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of the oxide scale. Besides, a strong oxidised state appears in the Mn 2p spectra. As
shown in Fig. 5e, the Mn 2p3/2 spectral peak with binding energy in the range of 640.4
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eV (at the surface) and 642.2 eV (at the sub-surface) corresponds to Mn oxides [26] and
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the metallic state with the binding energy of 638.8 eV (Mn 2p3/2) appears at the deeper
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regions.
Fig. 5. XPS survey and high-energy resolution spectra obtained at different sputtering depths of “oxidised surface II” (oxidised in a low vacuum of 10-1 Torr): (a) survey 13
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spectra, (b) Si 2p spectra, (c) Cr 2p spectra, (d) Fe 2p spectra and (e) Mn 2p spectra.
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Fig. 6. Deconvolution of XPS high-resolution spectra obtained after different depths of sputtering into oxide scale of “oxidised surface II” (oxidised in a low vacuum of
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10-1 Torr): (a-c) Fe 2p3/2 spectra at sputtering depths of 0 nm, 5.9 nm and 236 nm,
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respectively. (d-e) Mn 2p3/2 spectra at sputtering depths of 0 nm and 5.9 nm.
Figs. 6a-c illustrate the deconvolution of Fe 2p3/2 spectra obtained at depths of 0 nm,
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5.9 nm and 236 nm of “oxidised surface II”, respectively. The Fe 2p3/2 signal at the outmost surface (Fig. 6a) shows the presence of two components: trivalent metal ions (Fe3+) with a binding energy of 711.2 eV and bivalent metal ions (Fe2+) with a binding energy of 709.6 eV [26, 29]. The relative peak height indicates that Fe3+ is the primary 14
iron oxidized species at the surface region of the oxide scale, which may be due to the formation of MnFe2O4 or Fe2O3. As sputtering depth increases to 5.9 nm, the Fe2+ is the primary iron oxidized species (Fig. 6b), which should be attributed to the formation of (Fe, Cr)3O4 or FeO. After sputtering of 236 nm, only the Fe2+ component appears,
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indicating that Fe2+ is the only iron oxidized species at the deep region of the oxide scale. Moreover, the Mn 2p3/2 spectra obtained at the outmost surface (Fig. 6d) is fitted
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with two components. The component at 640.4 eV may belong to (Mn, Cr)3O4 or MnO, and the second component at 641.5 eV refers to a possible formation of MnFe2O4 or
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Mn2O3 [21, 26, 30]. After deep sputtering of 5.9 nm, the intensity component at 641.5
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eV becomes the primary Mn oxidised component (Fig. 6e).
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3.2 Cross-sectional analyses of “oxidised surface I” and “oxidised surface II”
Fig. 7. (a) High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image of the cross-section of “oxidised surface I”. (b) EDS line 15
scan profile across the oxide scale indicated by the blue line in (a). (c) Elemental maps of the area indicated by the blue rectangle in (a).
Fig. 7a shows the high-angle annular dark-field scanning transmission electron
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microscopy (HAADF-STEM) image of the cross-section of “oxidised surface I”. The thickness of the oxide scale is not uniform and the average thickness is approximately
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30 nm corresponding to the XPS result (Fig. 2b). Figs. 7b and c show the line scan profile and elemental maps of the oxide scale indicated by the blue line and rectangle
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in Fig. 7a, respectively. The results indicate that the oxide scale is composed of an inner
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Fe- and Cr-rich layer and an outer Fe-rich layer. High resolution transmission electron
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microscopy (HRTEM) image (Fig. 8) of the oxide scale formed on “oxidised surface I”
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shows that the interface of the matrix and oxide scale is atomically flat. Combining the
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HRTEM analysis (Fig. 8b), EDS results (Fig. 7b) and XPS results (Figs. 4e-f), the outer and inner layers of the oxide scale were determined as face-centered cubic (fcc)-
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structured Fe3O4 and FeCr2O4 spinel, respectively.
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Fig. 8. (a) High-resolution transmission electron microscopy (HRTEM) image of the
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oxide scale formed on “oxidised surface I”. (b) Enlarged image and (c) fast Fourier
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transform (FTT) of the region from rectangle shown on (a).
Fig. 9. (a) HAADF-STEM image of the cross-section of “oxidised surface II”. (b) EDS line scan profile across the oxide scale indicated by the blue line on (a). (c) Elemental maps of the oxide scale indicated by the blue rectangle on (a). (d) EDS line 17
scan profile across the oxide particle in internal oxidation zone (IOZ) indicated by the red line on (a).
Fig. 9a shows the cross-sectional morphology of “oxidized surface II”. It is indicated
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that the thickness of oxide scale is approximately 500 nm, in good agreement with the XPS result in Fig. 2c. Figs. 9b and c show the EDS line scan profile and elemental maps
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across the oxide scale indicated by the blue line and rectangle in Fig. 9a. The oxide
scale is comprised of two layers: the outer layer is composed of Fe, Mn and O, and the
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inner layer contains Mn, Cr, Fe and O. At the oxide scale/matrix interface, a Ni- and
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Mo-rich layer can be observed. Additionally, Si-rich nodules were observed in the inner
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layer. Fig. 10a presents the TEM bright-field image of the same area of the oxide scale
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as shown in Fig. 9a. The EDS results (Figs. 9b-c) and HRTEM image (Fig. 10b) indicate
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the presence of fcc-structured MnFe2O4 in the outer layer of the oxide scale. Moreover, the EDS results (Figs. 9b-c) and electron diffraction patterns (Figs. 10c-d) of the areas
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C and D as indicated in Fig. 10a demonstrate that the inner oxide layer is fcc-structured
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FeCr2O4 spinel with a lattice parameter of approximately 0.84, which is in agreement with the XPS results and Ref [31]. HRTEM image of the oxide scale/matrix interface (Fig. 10e) shows a neatly arranged atomic interface, indicating that the inner oxide layer
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is well bonded to the Ni- and Mo-rich layer. Beneath the oxide scale, an internal oxidation zone (IOZ) with several small oxide particles exist. These oxide particles were confirmed as MnCr2O4 according to the EDS line scan profile in Fig. 9d and Ref [32]. 18
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Fig. 10. (a) TEM bright-field image of the cross-section of “oxidized surface II” (the area indicated by green rectangle corresponds to the area indicated by the blue
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rectangle in Fig. 9a). (b, e) HRTEM images of areas B and E indicated by red
rectangle in (a). (c-d) Selected-area electron diffraction (SAED) images of areas C and D indicated by red circles in (a) from the [011] zone axis.
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3.3 Bonding characteristics of oxidised surfaces under different vacuum conditions
Figs. 11a-b present the bonding interfaces after hot-compression bonding under two different vacuum conditions. Before hot-compression bonding, “oxidised surface I” and
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“oxidised surface II” formed during initial heating and holding in a vacuum of 10-4 Torr
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and 10-1 Torr, respectively. After hot-compression bonding, a well-bonded interface was obtained in a high vacuum of 10-4 Torr and the original bonding interface transformed
into grain boundaries (Fig. 11a). While for the bonding interface obtained in a low
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vacuum of 10-1 Torr, discontinuous interfacial layers were observed at the bonding
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interface (Fig. 11b), which impeded the bonding of the two contacting surfaces. The
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EDS result (Fig. 11c) indicates that the interfacial layers are oxides.
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Fig. 11 Bonding interfaces obtained under different vacuum conditions. (a) Bonding interface of two contacting surfaces of “oxidised surface I” under high vacuum of 10-4 Torr, (b) Bonding interface of two contacting surfaces of “oxidised surface II” under
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low vacuum of 10-1 Torr. (c) EDS spectra obtained from point A in (b).
Furthermore, TEM experiments were carried out to further understand the structure
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and configuration of the interfacial oxides shown in Fig. 11b. HAADF-STEM image
(Fig. 12a) of the bonding interface in case of “oxidised surface II” shows that spheroidal,
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rod-like or irregular shaped interfacial oxides discontinuously distributed along the
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interface. EDS results indicate that the interfacial oxide is mainly composed of Mn, Cr,
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and O with the atomic ratio of approximately 1:2:4 (Figs. 12b-c). Combining with the
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HRTEM image (Fig. 12d), the interfacial oxide is determined to be fcc-structured
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MnCr2O4 with the lattice parameter of approximately 0.84 nm, which is consistent with the results reported in Ref. [32]. In addition, Elemental maps (Fig. 12b) also show the
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Si enrichment inside the interfacial oxides, indicating that Si-rich nodules in the scale
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of “oxidised surface II” (Fig. 9) remained at the bonding interface after the hot-
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Fig. 12. (a) HAADF-STEM image of the bonding interface formed in a low vacuum of 10-1 Torr with “oxidised surface II”. (b) Elemental maps of the boxed region and
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(c) line scan profile corresponding to the red line. (d) HRTEM image of interfacial
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oxide MnCr2O4 at the bonding interface.
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4. Discussion
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The thickness, composition and surface morphology of the oxide scales depend
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strongly on the oxidizing conditions. In particular, under vacuum conditions, the oxidation rate is mainly controlled by the impingement rate of oxygen molecules on the
J O2 [33]. According to the Hertz-Langmuir equation,
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specimen surface J O2
PO2
(2 M O2 RT )
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2
, where PO2
is the oxygen partial pressure in the partial
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vacuum, M O2 is the molecular mass of oxygen, R is the gas constant, T is the absolute temperature, and α is the adherence coefficient. Evidently, J O2 is positively related to the oxygen partial pressure, thus the oxidation rate in a high vacuum is slower than that in a low vacuum. Besides, the high vacuum (i.e. low oxygen partial pressure) 22
contributes to the formation of the initial protective Cr-rich layer due to selective oxidation [34]. The presence of such a Cr-rich layer tends to decrease the growth rate of oxide scale, acting as a barrier against the diffusion of metal and oxygen ions. Therefore, the final oxide scale formed in the high vacuum is very thin with a nanoscale
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thickness of 30 nm and it consists of outer Fe3O4 layer and inner Fe- and Cr-rich layer. While in the low vacuum (10-1 Torr), the oxide scale grows rapidly owing to the high
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J O2 , which makes this case different. Not only is the oxide scale much thicker (500
nm), but also its composition changes. Here, MnFe2O4 was observed in the outer layer
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rather than Fe3O4 or MnCr2O4 as identified in previous studies [24, 35]. The inner layer
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is primarily composed of Cr-rich FeCr2O4 spinel with a small fraction of MnFe2O4. The
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distribution of Fe, Mn and Cr in this oxide scale could be related to their different
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affinities for oxygen and diffusion rates in oxide scale. As the oxygen affinity of
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chromium is stronger than those of iron and other elements in alloy, oxygen is more inclined to react with Cr, which leads to initial formation of a protective inner chromia
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scale. When the Cr supply becomes insufficient to support the continuously growing
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chromia scale. Chromia is then gradually enriched in Fe and eventually converted in less protective FeCr2O4 spinel [24]. It results in fast outward diffusion of metal ions and inward diffusion of oxygen. As the diffusion rates of Mn and Fe are faster than those of
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Cr, Ni and Mo in the FeCr2O4 spinel [36, 37], the high concentrations of Mn and Fe diffuse from the matrix to surface through inner oxide layer, thus leading to an outward growing MnFe2O4 oxide. In addition, Ni and Mo enrichments at the oxide scale/matrix interface were observed 23
under the oxide scale formed in the low vacuum. Previous studies have reported the similar Ni enrichment at the oxide scale/matrix interface [19, 38]. Wang et al. [19] reported the Mo enrichment in the inner oxide scale. While to the best of our knowledge, Mo enrichment at the oxide scale/matrix interface has rarely been mentioned. However,
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in the present study, both Ni and Mo enrichments at the oxide scale/matrix interface were observed. This is because Ni and Mo are the least likely elements to be oxidized
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owing to their lowest oxygen affinities among all the elements in austenite stainless steel [39]. Additionally, the diffusion rates of Ni and Mo are slow compared with Mn,
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Fe and Cr [40]. As a result, Cr with the highest oxygen affinity is oxidized to form the
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inner FeCr2O4 scale, meanwhile, Fe and Mn diffuse through inner oxide layer to form
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between the oxide scale and matrix.
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the outer Fe- and Mn-rich oxide layer. Further, Ni and Mo are left at the interface
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Fig. 13. Property diagrams versus oxygen activity for 316LN stainless steel at
1200 °C calculated using Thermo-Calc with TCFE8 database: (a) Volume fraction of phases versus oxygen activity, (b) Atomic fraction of all elements in spinel oxide versus oxygen activity.
24
It is noteworthy that the final interfacial oxides consist of MnCr2O4 spinel along with small amounts of Si-rich nodules (Fig. 12) after bonding in a low vacuum of 10-1 Torr. While the initial surface oxide scale on the contacting surface (i.e. “oxidised surface II”) is a duplex layer structure composed of mixed oxides such as FeCr2O4, MnFe2O4
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and Si-rich oxide. This means that a transition from the surface oxide scale to interfacial oxides occurs during the bonding process. In order to discuss this conversion,
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thermodynamic calculations were performed using Thermo-Calc software with TCFE8 data base [41, 42]. The results are presented in Fig. 13. Once two “oxidised surfaces II”
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contact with each other, the surface oxide scales are broken up under the compression
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deformation and remain on the bonding interface (Fig. 11b). As the broken oxide scales
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act as oxygen source of the interface during the bonding process, the interfacial oxygen
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activity is controlled by the dissociation oxygen activity of the broken oxide scales.
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Since the oxygen activity in the matrix is lower than that at the interface, the broken oxide scales on the bonding interface are unstable and the dissolved oxygen ions will
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diffuse toward the matrix. Therefore, the oxygen activity at the bonding interface
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decreases continuously and thus the composition of the spinel oxide scale is transformed from Fe–Cr and Fe–Mn spinel into Mn–Cr spinel according to Figs. 13ab. Thus after cooling, only final MnCr2O4 along with small amounts of Si-rich nodules
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remain at the bonding interface. Only if subsequent holding treatment on this bonding joint at 1200 °C is carried out, the final interfacial oxides MnCr2O4 will decompose, as reported in our previous study [43]. By contrast, the bonding interface formed in the high vacuum of 10-4 Torr is well bonded and no interfacial oxides are observed at the 25
bonding interface. This is because the initial oxide scale on the contacting surface (i.e. “oxidised surface I”) is mainly composed of Fe and Cr oxides and it has a nanoscale thickness. Similarly, once two “oxidised surfaces I” contact with each other, the Fe- and Cr-rich oxide scales are broken up into isolated particles distributed on the bonding
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interface due to deformation. The nanoscale Fe- and Cr-rich oxide particles are unstable and decompose immediately during hot-compression bonding. Thus no interfacial
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oxides exist on the bonding interface without further holding treatments.
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5. Conclusions
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The surface chemistry and structure of initial surface oxide scales formed on 316LN
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stainless steel and the interfacial characteristics after hot-compression bonding under
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different vacuum conditions were investigated using a combination of XPS, TEM, SEM
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and thermodynamic calculations. We report the following findings: (1) The oxide scale of as-received steel consists of binary oxide of Fe and Cr with a
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thickness of approximately 4 nm. After holding in a high vacuum of 10-4 Torr, an oxide
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scale with a thickness of 30 nm is observed. While in a low vacuum of 10-1 Torr, the oxide scale is a duplex structure with a thickness of approximately 500 nm.
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(2) The initial oxide scale formed in a high vacuum of 10-4 Torr is composed of outer
Fe3O4 and inner Fe- and Cr-rich layer. After hot-compression bonding, the bonding interface is well bonded and no interfacial oxides are observed on the interface, owing to the nanoscale thickness and instability of the initial oxide scale. (3) The initial oxide scale formed in a low vacuum of 10-1 Torr is composed of outer 26
MnFe2O4 layer and inner layer consisting of mixed FeCr2O4 and MnFe2O4 along with small amounts of Si-rich nodules. Nickel and molybdenum enrichments are observed at the oxide scale/matrix interface. The thick complex oxide scale on the contacting surfaces transforms into MnCr2O4 after hot-compression bonding, which is resulted
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from the decrease of oxygen activity at the bonding interface. (4) Different vacuum conditions result in different thicknesses and types of oxide
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scales. A higher vacuum leads to a thinner and unstable oxide scale, which decomposes
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easily during the hot-compression bonding process, leading to a well-bonded interface.
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Acknowledgements
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The authors acknowledge the financial support from National key Research and
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Development program (Grant No. 2016YFB0300401), National Natural Science
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Foundation of China (Grant Nos. U1508215, 51774265) and key Program of the
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S0010-938X(18)31547-6 https://doi.org/10.1016/j.corsci.2018.11.001 CS 7757
To appear in: Received date: Revised date: Accepted date:
17 September 2018 5 November 2018 10 November 2018
Please cite this article as: Xie B, Sun M, Xu B, Wang C, Jiang H, Li D, Li Y, Oxidation of stainless steel in vacuum and evolution of surface oxide scales during hot-compression bonding, Corrosion Science (2018), https://doi.org/10.1016/j.corsci.2018.11.001 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Oxidation of stainless steel in vacuum and evolution of surface oxide scales during hot-compression bonding Bijun Xiea,b, Mingyue Suna,d,*, Bin Xua,d, Chunyang Wanga,b, Haiyang Jiangc,
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Shenyang National Laboratory for Materials Science, Institute of Metal Research,
b
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Chinese Academy of Sciences, Shenyang 110016, China
c
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Dianzhong Lia, Yiyi Lia
University of Chinese Academy of Sciences, Beijing 100049, China
School of Metallurgy and Materials Engineering, Jiangsu University of Science and
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Key Laboratory of Nuclear Materials and Safety Assessment, Institute of Metal
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d
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Technology, Zhenjiang 212003, China
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Research, Chinese Academy of Sciences, Shenyang 110016, China
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*Corresponding author. Tel.: +86 24 83978839; E-mail: [email protected]
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Highlights
The oxidation behaviors of austenite stainless steel in vacuum were investigated.
Thick oxide scale with outer MnFe2O4 and inner FeCr2O4 formed in a low vacuum.
Complex surface oxide scale transformed to MnCr2O4 after hot compression
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bonding.
Oxide scale with outer Fe3O4 and inner (Fe, Cr)-rich layer formed in a high vacuum.
Nanoscale-thick surface oxide scale decomposed during hot compression bonding.
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Abstract
The oxidation behaviour of austenite stainless steel in vacuum and the evolution of oxide scales during vacuum hot-compression bonding were investigated combining
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XPS and TEM. The results indicate that thick duplex oxide scale with outer MnFe2O4 and inner FeCr2O4 layer was formed in a low vacuum and the complex oxide scale
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transformed into interfacial oxides MnCr2O4 remaining at the bonding interface after
hot-compression bonding at 1200 °C. While in a high vacuum, nanoscale-thick oxide scale with outer Fe3O4 and inner Fe- and Cr-rich layer was observed and it decomposed
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during the subsequent hot-compression bonding procedure, leading to a well-bonded
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interface.
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Keywords: Stainless steel; Vacuum oxidation; XPS; TEM; Hot-compression bonding
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1. Introduction
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Since stainless steel clad plates simultaneously exhibit good mechanical properties and corrosion resistance, they have attracted increasing attention for various
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engineering applications, e.g. heat exchangers, nuclear power equipment, transferring pipes, vessels, and automobiles [1-3]. Stainless steel clad plates are layered structures
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that are generally fabricated by bonding stainless steel plates to other metal plates, such as carbon steel, low-alloy steel, titanium alloy, or aluminium alloy [4-7]. Different technologies are used to manufacture stainless steel clad plates, e.g. diffusion bonding, explosion welding, and roll bonding [4, 8-10]. Among all technologies, hot roll bonding is the most effective and suitable manufacturing procedure for clad plates during which 2
the metal plates are joined under high temperatures [11]. However, regarding the conventional hot roll bonding technique, as metal surfaces significantly react with atmospheric oxygen, oxide scales form on the contacting surfaces during heating. This deteriorates the bonding quality of the clad plates. Therefore, vacuum rolling cladding
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(VRC) is applied to obtain better bonding effect during which the contacting surfaces of the metal plates are cleaned by mechanical polishing and chemical etching to remove
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barriers, and then bonded in a vacuum. However the fact is that oxidation at the bonding
interface is still inevitable even under vacuum conditions. For example, Xie et al. [12]
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showed that the Mn–Si–O oxide particles were formed at the clad interface during VRC
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process in a vacuum of 10-1 Pa. Zhu et al. [13] reported the formation of interfacial
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oxide at 316L stainless steel cladding interface and indicated that the interfacial oxide
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is the limiting factor for interface bonding. Actually, these interfacial oxides at the clad
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interface are determined by the initial surface oxidation of the contacting surfaces during initial holding process under high temperatures and vacuum conditions before
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rolling cladding. Therefore, it’s essential to investigate the initial high-temperature
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oxidation of the contacting surfaces under vacuum conditions in order to control and eliminate interfacial oxides to acquire better bonding effect. This is particularly necessary for stainless steel clad plates with compact and stable oxide scales on the
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stainless steel surface. The oxidation behaviour of stainless steel in super critical water, water vapour, steam and ambient air at high temperature has been extensively studied [14-18]. In all these cases, it has been comprehensively reported that the oxide scale has a duplex structure, 3
composed of Fe and Cr oxides. The outer layer of the scale is predominantly composed of Fe-rich oxide, whereas the inner layer consists of Cr-rich oxide. Furthermore, Li et al. [19] reported the early oxidation mechanism of stainless steel TP347H in supercritical water. Matthews et al. [20] investigated the intergranular oxidation of 316L
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stainless steel in the PWR water environment and suggested that an epitaxial relationship existed between the metal and the inner Cr-rich oxide. Swaminathan et al.
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[21, 22] reported that the oxide scale formed in ambient air was mainly composed of
mixed oxides such as FeCr2O4 and MnCr2O4 along with binary oxides with Fe, Cr, and
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Mn. Col and Cheng et al. [23, 24] focused on the breakaway oxidation behaviour of
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stainless steel at high temperature. It is reported that the breakaway oxidation behaviour
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was significantly influenced by the microstructure and the composition of the oxide
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scale and the occurrence of breakaway oxidation was related to a local conversion of
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Cr2O3 in less protective FeCr2O4. Ostwald et al. [25] found that the oxide scale formed on stainless steel in H2–2.5%H2O atmosphere was protective and contained high Cr,
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Mn and Si contents, while in N2-20.5%O2 atmosphere, the oxide scale was Fe- and Ni-
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containing oxides. Nevertheless, to the best of our knowledge, systematic investigations on the high-temperature oxidation of stainless steel under vacuum conditions are scarce. The objective of this work is to investigate the oxidation behaviour of austenite
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stainless steel in vacuum and the evolution of oxide scales during vacuum hotcompression bonding. By using X-ray photoelectron spectroscopy (XPS) and transmission electron microscopy (TEM), we systematically studied the surface chemistry and structure of oxide scales formed on the surface of 316LN stainless steel 4
under two different vacuum conditions (10-1 Torr and 10-4 Torr). Moreover, the evolution of the surface oxide scales during hot-compression bonding in corresponding vacuum environments was investigated. The results show that the initial surface complex oxide scale transformed into an interfacial MnCr2O4 layer during hot-
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compression bonding in a low vacuum, whereas the nanoscale-thick surface oxide scale formed in a high vacuum has decomposed after hot-compression bonding. The
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systematic investigation of the surface oxidation and evolution of oxide scales during
vacuum hot-compression bonding provide new insight into high-temperature oxidation
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in vacuum and may have implications in manufacturing of high-quality bonding joints.
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2. Experimental method
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Austenite stainless steel (316LN) with a chemical composition (in wt.%) of 0.012 C,
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0.12 N, 0.34 Si, 0.003 S, 0.018 P, 12.83 Ni, 1.41 Mn, 2.18 Mo, 16.93 Cr, 0.016 Al, 0.04 Cu, 0.059 Nb and balance Fe was used as raw material. Square samples with the
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dimensions of 8*8*6 mm were cut from the raw material. The test surfaces were wet-
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ground and buff-polished with diamond paste. The buff-polished surfaces are named as “native surface”. Two groups of square samples were heated in the Gleeble-3800
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thermal simulator under different vacuum conditions, i.e. 10-4 Torr and 10-1 Torr, respectively. During the heat treatment, the samples were first heated to 1200 °C and maintained for 5 min. Afterwards, the samples were cooled to reach 30 °C and then removed from the vacuum environment. The surfaces of the treated samples are called “oxidised surface I” and “oxidised surface II” for two vacuum environments of 10-4 5
Torr and 10-1 Torr, respectively. After this treatment, the “native surface”, “oxidised surface I”, and “oxidised surface II” were investigated to understand the oxidation of the contacting surfaces under different vacuum conditions using XPS and TEM. Furthermore, the hot-compression bonding experiments (Fig. 1a) were performed in the
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Gleeble-3800 with vacuum degree of 10-4 Torr and 10-1 Torr at the strain rate of 0.1. The interfacial characteristics of two kinds of bonding interfaces were analysed by
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scanning electron microscope (SEM) and TEM.
The XPS measurements were performed using a PHI Quantera II equipped with a
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monochromatic Al Kα (1486.6 eV) X-ray source. To determine the chemistry of the
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oxide scales formed on the “native surface”, “oxidised surface I” and “oxidised surface
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II”, XPS survey scans and high-energy resolution spectra of characteristic peaks of
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elements such as Fe 2p, Cr 2p, Mn 2p, Si 2p, C 1s, and O 1s were acquired. In order to
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determine the elemental distribution in depth from the surface, i.e. the oxide chemistry in the sub-surface zones, the samples were sputtered with Ar+ ions with an energy of 2
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keV. The elemental spectra were fitted using Casa XPS software with Gaussian–
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Lorentzian line shapes and the Shirley-type background.
Fig. 1. (a) Schematic of hot-compression bonding process of two specimens under different vacuum conditions. (b) Preparation of transmission electron microscopy 6
(TEM) sample by focused ion beam (FIB) and (c) thinning of the sample lamella with the oxide scale on it using FIB.
The TEM sample containing the oxide scale was prepared and extracted using a
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focused ion beam (FIB) (FEI Helios NanoLab 650 DualBeam) with the in-situ loft-out (INLO) method, as shown in Fig. 1b. After which, the sample was further thinned to
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electron transparency by ion milling (Fig. 1c). Detailed structural and elemental
characterizations of the oxide scales and interfacial oxides were conducted using an FEI
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Tecnai G2 F20 transmission electron microscope attached with a high-angle annular
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dark-field detector and an EDS system with an acceleration voltage of 200 kV.
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3. Results
3.1 XPS analyses of oxide scales formed on “native surface”, “oxidised surface I”, and “oxidised surface II”
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3.1.1 Depth analyses
Fig. 2. Variations in chemical composition with respect to depth: (a) “native surface”,
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(b) “oxidised surface I” (treated in a high vacuum of 10-4 Torr), (c) “oxidised surface II” (treated in a low vacuum of 10-1 Torr).
The variations in the chemical composition with depth were obtained on three types of oxidised surfaces using XPS (Fig. 2). The XPS oxygen profile in Fig. 2a shows that 8
the oxide scale on the “native surface” is approximately 4 nm thick and there is a decreasing oxygen gradient, after which the composition of the matrix starts. Fe and Cr enrichment is visible at the surface (top 4 nm). As for the “oxidised surface I” (treated in a high vacuum of 10-4 Torr), the oxide scale thickness is approximately 30 nm and a
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decreasing oxygen gradient over ~50 nm is observed (Fig. 2b). It can be seen that the Fe and Cr enrichment is obvious throughout the oxide scale. Fig. 2c shows an overview
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composition of the top 3.6 μm of “oxidised surface II” (treated in a low vacuum of 10 Torr) and this oxide scale thickness is approximately 500 nm. Further, an internal
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oxidation zone (IOZ) of approximately 2.5 μm exists, after which the composition of
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the matrix starts. The initial Fe and Mn enrichment around 50 nm in depth in Fig. 2c
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indicates that Fe- and Mn-rich oxides are likely formed on the top surface. The obvious
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Cr enrichment at the surface (top 500 nm) demonstrates the presence of Cr-rich oxides.
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Besides, the Si enrichment is also observed in the top 500 nm.
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3.1.2 Surface chemistry
Fig. 3. XPS survey and high-energy resolution spectra obtained at different sputtering
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depths of “native surface”: (a) survey spectra, (b) Cr 2p spectra and (c) Fe 2p spectra.
Fig. 3 presents the XPS survey scan and high-resolution spectra of “native surface”.
The depth profiles were obtained for different sputtering-time intervals to assess the
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chemical state of elements at different depths of the oxide scales. The outmost surface spectra Cr 2p3/2 and Fe 2p3/2 (Figs. 3b and c) have binding energies of 576.1 eV and 710.37 eV, corresponding to Cr oxide and Fe oxide, respectively [26]. This confirms that binary oxides of Cr and Fe are formed on the initial “native surface”. Additionally, 10
tiny shoulders at 573.46 eV (Cr 2p3/2), 706.33 eV (Fe 2p3/2) and 638.8 eV (Mn 2p3/2) indicate their metallic state [26] and these metallic spectral peaks increase with increasing sputter depth. After sputtering down to 5 nm, only metallic spectral peaks appear (Figs. 3a-c), which indicates that the thickness of the oxide scale formed on the
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“native surface” is below 5 nm.
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Fig. 4. XPS survey and high-resolution spectra obtained at different sputtering depths
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of “oxidised surface I” (oxidised in a high vacuum of 10-4 Torr): (a) survey spectra, (b) Si 2p spectra, (c) Cr 2p spectra, (d) Fe 2p spectra, (e) Curve fitted Fe 2p3/2 spectra
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obtained at sputtering depth of 0 nm, (f) Deconvolution of Fe 2p3/2 spectra obtained at sputtering depth of 12 nm.
In-depth distributions of the XPS survey and high-resolution spectra of Si 2p, Cr 2p and Fe 2p of “oxidised surface I” are presented in Fig. 4. According to the survey 11
spectra (Fig. 4a), The Mn spectral peak at 638.6 eV (Mn 2p3/2) was only observed in the deeper region, indicating its metallic state. A Si 2p spectral peak was also detected, and the Si 2p3/2 spectral peak (103.4 eV) [26] only appeared at the surface regions, corresponding to the oxidized Si. No spectral peaks were observed at deeper regions
IP T
owing to the low Si content in the matrix (Fig. 4b). Regarding depth-resolved Cr 2p spectra, the binding energy of 576.6 eV for Cr 2p3/2 spectrum in Fig. 4c corresponds to
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Cr oxide [26, 27]. Similarly, after sputtering down for a certain depth, a metallic Cr peak appeared at 573.8 eV (Cr 2p3/2) [27]. The depth-resolved Fe 2p3/2 spectra in Fig.
U
4d positioned the binding energy at 710.5 eV (at the surface), indicating that Fe exists
N
in oxidized state of Fe3O4 at the surface [26] (Fig. 4e). With increasing sputtering depth,
A
Fe 2p3/2 peak shifted towards a smaller binding energy and broadened (at the sub-
M
surface), suggesting a mixed Fe state. Fig. 4f illustrates the deconvolution of Fe 2p 3/2
ED
spectrum obtained after 12 nm sputtering into the oxide scale formed on “oxidised surface I”. Three components were necessitated to fit the Fe 2p3/2 transition. The higher
PT
binding energy component at 710.5 eV is attributed to Fe3O4 [26], and the lower binding
CC E
energy component at 709 eV is due to Fe2+ which may be from FeO or (Fe,Cr)3O4 oxides [26, 28]. The binding energy component at 706.8 eV belongs to metallic Fe [26], and the intensity of the metallic spectral peak increases with increasing sputter depth
A
(Fig. 4d).
Fig. 5 shows the XPS survey and high-resolution (Si 2p, Cr 2p, Fe 2p and Mn 2p) spectra obtained on the “oxidised surface II” at different depths. The in-depth distributions of Si 2p and Cr 2p spectra (Figs. 5b and c) have similar variations as that 12
of the “oxidised surface I” (Figs. 4b and c). As for Fe 2p spectra (Fig. 5d), the binding energy of Fe 2p3/2 spectra range from 710.7 eV (at the surface) to 709.4 eV (at the subsurface) corresponds to Fe oxides [26]. This oxide spectral peak gradually weakens with increasing sputter depth towards the matrix and the metallic Fe peak at 706.9 eV
IP T
(Fe 2p3/2) appears (Fig. 5d) [26]. It is noteworthy that the intensity of Fe 2p3/2 peak is weak at the depth of 236 nm, which indicates that Fe oxide content is low at this region
SC R
of the oxide scale. Besides, a strong oxidised state appears in the Mn 2p spectra. As
shown in Fig. 5e, the Mn 2p3/2 spectral peak with binding energy in the range of 640.4
U
eV (at the surface) and 642.2 eV (at the sub-surface) corresponds to Mn oxides [26] and
N
the metallic state with the binding energy of 638.8 eV (Mn 2p3/2) appears at the deeper
A
CC E
PT
ED
M
A
regions.
Fig. 5. XPS survey and high-energy resolution spectra obtained at different sputtering depths of “oxidised surface II” (oxidised in a low vacuum of 10-1 Torr): (a) survey 13
M
A
N
U
SC R
IP T
spectra, (b) Si 2p spectra, (c) Cr 2p spectra, (d) Fe 2p spectra and (e) Mn 2p spectra.
ED
Fig. 6. Deconvolution of XPS high-resolution spectra obtained after different depths of sputtering into oxide scale of “oxidised surface II” (oxidised in a low vacuum of
PT
10-1 Torr): (a-c) Fe 2p3/2 spectra at sputtering depths of 0 nm, 5.9 nm and 236 nm,
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respectively. (d-e) Mn 2p3/2 spectra at sputtering depths of 0 nm and 5.9 nm.
Figs. 6a-c illustrate the deconvolution of Fe 2p3/2 spectra obtained at depths of 0 nm,
A
5.9 nm and 236 nm of “oxidised surface II”, respectively. The Fe 2p3/2 signal at the outmost surface (Fig. 6a) shows the presence of two components: trivalent metal ions (Fe3+) with a binding energy of 711.2 eV and bivalent metal ions (Fe2+) with a binding energy of 709.6 eV [26, 29]. The relative peak height indicates that Fe3+ is the primary 14
iron oxidized species at the surface region of the oxide scale, which may be due to the formation of MnFe2O4 or Fe2O3. As sputtering depth increases to 5.9 nm, the Fe2+ is the primary iron oxidized species (Fig. 6b), which should be attributed to the formation of (Fe, Cr)3O4 or FeO. After sputtering of 236 nm, only the Fe2+ component appears,
IP T
indicating that Fe2+ is the only iron oxidized species at the deep region of the oxide scale. Moreover, the Mn 2p3/2 spectra obtained at the outmost surface (Fig. 6d) is fitted
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with two components. The component at 640.4 eV may belong to (Mn, Cr)3O4 or MnO, and the second component at 641.5 eV refers to a possible formation of MnFe2O4 or
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Mn2O3 [21, 26, 30]. After deep sputtering of 5.9 nm, the intensity component at 641.5
A
N
eV becomes the primary Mn oxidised component (Fig. 6e).
A
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PT
ED
M
3.2 Cross-sectional analyses of “oxidised surface I” and “oxidised surface II”
Fig. 7. (a) High-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) image of the cross-section of “oxidised surface I”. (b) EDS line 15
scan profile across the oxide scale indicated by the blue line in (a). (c) Elemental maps of the area indicated by the blue rectangle in (a).
Fig. 7a shows the high-angle annular dark-field scanning transmission electron
IP T
microscopy (HAADF-STEM) image of the cross-section of “oxidised surface I”. The thickness of the oxide scale is not uniform and the average thickness is approximately
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30 nm corresponding to the XPS result (Fig. 2b). Figs. 7b and c show the line scan profile and elemental maps of the oxide scale indicated by the blue line and rectangle
U
in Fig. 7a, respectively. The results indicate that the oxide scale is composed of an inner
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Fe- and Cr-rich layer and an outer Fe-rich layer. High resolution transmission electron
A
microscopy (HRTEM) image (Fig. 8) of the oxide scale formed on “oxidised surface I”
M
shows that the interface of the matrix and oxide scale is atomically flat. Combining the
ED
HRTEM analysis (Fig. 8b), EDS results (Fig. 7b) and XPS results (Figs. 4e-f), the outer and inner layers of the oxide scale were determined as face-centered cubic (fcc)-
A
CC E
PT
structured Fe3O4 and FeCr2O4 spinel, respectively.
16
IP T SC R U
N
Fig. 8. (a) High-resolution transmission electron microscopy (HRTEM) image of the
A
oxide scale formed on “oxidised surface I”. (b) Enlarged image and (c) fast Fourier
A
CC E
PT
ED
M
transform (FTT) of the region from rectangle shown on (a).
Fig. 9. (a) HAADF-STEM image of the cross-section of “oxidised surface II”. (b) EDS line scan profile across the oxide scale indicated by the blue line on (a). (c) Elemental maps of the oxide scale indicated by the blue rectangle on (a). (d) EDS line 17
scan profile across the oxide particle in internal oxidation zone (IOZ) indicated by the red line on (a).
Fig. 9a shows the cross-sectional morphology of “oxidized surface II”. It is indicated
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that the thickness of oxide scale is approximately 500 nm, in good agreement with the XPS result in Fig. 2c. Figs. 9b and c show the EDS line scan profile and elemental maps
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across the oxide scale indicated by the blue line and rectangle in Fig. 9a. The oxide
scale is comprised of two layers: the outer layer is composed of Fe, Mn and O, and the
U
inner layer contains Mn, Cr, Fe and O. At the oxide scale/matrix interface, a Ni- and
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Mo-rich layer can be observed. Additionally, Si-rich nodules were observed in the inner
A
layer. Fig. 10a presents the TEM bright-field image of the same area of the oxide scale
M
as shown in Fig. 9a. The EDS results (Figs. 9b-c) and HRTEM image (Fig. 10b) indicate
ED
the presence of fcc-structured MnFe2O4 in the outer layer of the oxide scale. Moreover, the EDS results (Figs. 9b-c) and electron diffraction patterns (Figs. 10c-d) of the areas
PT
C and D as indicated in Fig. 10a demonstrate that the inner oxide layer is fcc-structured
CC E
FeCr2O4 spinel with a lattice parameter of approximately 0.84, which is in agreement with the XPS results and Ref [31]. HRTEM image of the oxide scale/matrix interface (Fig. 10e) shows a neatly arranged atomic interface, indicating that the inner oxide layer
A
is well bonded to the Ni- and Mo-rich layer. Beneath the oxide scale, an internal oxidation zone (IOZ) with several small oxide particles exist. These oxide particles were confirmed as MnCr2O4 according to the EDS line scan profile in Fig. 9d and Ref [32]. 18
IP T SC R U N A M ED PT
CC E
Fig. 10. (a) TEM bright-field image of the cross-section of “oxidized surface II” (the area indicated by green rectangle corresponds to the area indicated by the blue
A
rectangle in Fig. 9a). (b, e) HRTEM images of areas B and E indicated by red
rectangle in (a). (c-d) Selected-area electron diffraction (SAED) images of areas C and D indicated by red circles in (a) from the [011] zone axis.
19
3.3 Bonding characteristics of oxidised surfaces under different vacuum conditions
Figs. 11a-b present the bonding interfaces after hot-compression bonding under two different vacuum conditions. Before hot-compression bonding, “oxidised surface I” and
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“oxidised surface II” formed during initial heating and holding in a vacuum of 10-4 Torr
SC R
and 10-1 Torr, respectively. After hot-compression bonding, a well-bonded interface was obtained in a high vacuum of 10-4 Torr and the original bonding interface transformed
into grain boundaries (Fig. 11a). While for the bonding interface obtained in a low
U
vacuum of 10-1 Torr, discontinuous interfacial layers were observed at the bonding
A
N
interface (Fig. 11b), which impeded the bonding of the two contacting surfaces. The
A
CC E
PT
ED
M
EDS result (Fig. 11c) indicates that the interfacial layers are oxides.
20
Fig. 11 Bonding interfaces obtained under different vacuum conditions. (a) Bonding interface of two contacting surfaces of “oxidised surface I” under high vacuum of 10-4 Torr, (b) Bonding interface of two contacting surfaces of “oxidised surface II” under
IP T
low vacuum of 10-1 Torr. (c) EDS spectra obtained from point A in (b).
Furthermore, TEM experiments were carried out to further understand the structure
SC R
and configuration of the interfacial oxides shown in Fig. 11b. HAADF-STEM image
(Fig. 12a) of the bonding interface in case of “oxidised surface II” shows that spheroidal,
U
rod-like or irregular shaped interfacial oxides discontinuously distributed along the
N
interface. EDS results indicate that the interfacial oxide is mainly composed of Mn, Cr,
A
and O with the atomic ratio of approximately 1:2:4 (Figs. 12b-c). Combining with the
M
HRTEM image (Fig. 12d), the interfacial oxide is determined to be fcc-structured
ED
MnCr2O4 with the lattice parameter of approximately 0.84 nm, which is consistent with the results reported in Ref. [32]. In addition, Elemental maps (Fig. 12b) also show the
PT
Si enrichment inside the interfacial oxides, indicating that Si-rich nodules in the scale
CC E
of “oxidised surface II” (Fig. 9) remained at the bonding interface after the hot-
A
compression bonding.
21
IP T SC R
Fig. 12. (a) HAADF-STEM image of the bonding interface formed in a low vacuum of 10-1 Torr with “oxidised surface II”. (b) Elemental maps of the boxed region and
N
U
(c) line scan profile corresponding to the red line. (d) HRTEM image of interfacial
A
oxide MnCr2O4 at the bonding interface.
M
4. Discussion
ED
The thickness, composition and surface morphology of the oxide scales depend
PT
strongly on the oxidizing conditions. In particular, under vacuum conditions, the oxidation rate is mainly controlled by the impingement rate of oxygen molecules on the
J O2 [33]. According to the Hertz-Langmuir equation,
CC E
specimen surface J O2
PO2
(2 M O2 RT )
1
2
, where PO2
is the oxygen partial pressure in the partial
A
vacuum, M O2 is the molecular mass of oxygen, R is the gas constant, T is the absolute temperature, and α is the adherence coefficient. Evidently, J O2 is positively related to the oxygen partial pressure, thus the oxidation rate in a high vacuum is slower than that in a low vacuum. Besides, the high vacuum (i.e. low oxygen partial pressure) 22
contributes to the formation of the initial protective Cr-rich layer due to selective oxidation [34]. The presence of such a Cr-rich layer tends to decrease the growth rate of oxide scale, acting as a barrier against the diffusion of metal and oxygen ions. Therefore, the final oxide scale formed in the high vacuum is very thin with a nanoscale
IP T
thickness of 30 nm and it consists of outer Fe3O4 layer and inner Fe- and Cr-rich layer. While in the low vacuum (10-1 Torr), the oxide scale grows rapidly owing to the high
SC R
J O2 , which makes this case different. Not only is the oxide scale much thicker (500
nm), but also its composition changes. Here, MnFe2O4 was observed in the outer layer
U
rather than Fe3O4 or MnCr2O4 as identified in previous studies [24, 35]. The inner layer
N
is primarily composed of Cr-rich FeCr2O4 spinel with a small fraction of MnFe2O4. The
A
distribution of Fe, Mn and Cr in this oxide scale could be related to their different
M
affinities for oxygen and diffusion rates in oxide scale. As the oxygen affinity of
ED
chromium is stronger than those of iron and other elements in alloy, oxygen is more inclined to react with Cr, which leads to initial formation of a protective inner chromia
PT
scale. When the Cr supply becomes insufficient to support the continuously growing
CC E
chromia scale. Chromia is then gradually enriched in Fe and eventually converted in less protective FeCr2O4 spinel [24]. It results in fast outward diffusion of metal ions and inward diffusion of oxygen. As the diffusion rates of Mn and Fe are faster than those of
A
Cr, Ni and Mo in the FeCr2O4 spinel [36, 37], the high concentrations of Mn and Fe diffuse from the matrix to surface through inner oxide layer, thus leading to an outward growing MnFe2O4 oxide. In addition, Ni and Mo enrichments at the oxide scale/matrix interface were observed 23
under the oxide scale formed in the low vacuum. Previous studies have reported the similar Ni enrichment at the oxide scale/matrix interface [19, 38]. Wang et al. [19] reported the Mo enrichment in the inner oxide scale. While to the best of our knowledge, Mo enrichment at the oxide scale/matrix interface has rarely been mentioned. However,
IP T
in the present study, both Ni and Mo enrichments at the oxide scale/matrix interface were observed. This is because Ni and Mo are the least likely elements to be oxidized
SC R
owing to their lowest oxygen affinities among all the elements in austenite stainless steel [39]. Additionally, the diffusion rates of Ni and Mo are slow compared with Mn,
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Fe and Cr [40]. As a result, Cr with the highest oxygen affinity is oxidized to form the
N
inner FeCr2O4 scale, meanwhile, Fe and Mn diffuse through inner oxide layer to form
CC E
PT
ED
M
between the oxide scale and matrix.
A
the outer Fe- and Mn-rich oxide layer. Further, Ni and Mo are left at the interface
A
Fig. 13. Property diagrams versus oxygen activity for 316LN stainless steel at
1200 °C calculated using Thermo-Calc with TCFE8 database: (a) Volume fraction of phases versus oxygen activity, (b) Atomic fraction of all elements in spinel oxide versus oxygen activity.
24
It is noteworthy that the final interfacial oxides consist of MnCr2O4 spinel along with small amounts of Si-rich nodules (Fig. 12) after bonding in a low vacuum of 10-1 Torr. While the initial surface oxide scale on the contacting surface (i.e. “oxidised surface II”) is a duplex layer structure composed of mixed oxides such as FeCr2O4, MnFe2O4
IP T
and Si-rich oxide. This means that a transition from the surface oxide scale to interfacial oxides occurs during the bonding process. In order to discuss this conversion,
SC R
thermodynamic calculations were performed using Thermo-Calc software with TCFE8 data base [41, 42]. The results are presented in Fig. 13. Once two “oxidised surfaces II”
U
contact with each other, the surface oxide scales are broken up under the compression
N
deformation and remain on the bonding interface (Fig. 11b). As the broken oxide scales
A
act as oxygen source of the interface during the bonding process, the interfacial oxygen
M
activity is controlled by the dissociation oxygen activity of the broken oxide scales.
ED
Since the oxygen activity in the matrix is lower than that at the interface, the broken oxide scales on the bonding interface are unstable and the dissolved oxygen ions will
PT
diffuse toward the matrix. Therefore, the oxygen activity at the bonding interface
CC E
decreases continuously and thus the composition of the spinel oxide scale is transformed from Fe–Cr and Fe–Mn spinel into Mn–Cr spinel according to Figs. 13ab. Thus after cooling, only final MnCr2O4 along with small amounts of Si-rich nodules
A
remain at the bonding interface. Only if subsequent holding treatment on this bonding joint at 1200 °C is carried out, the final interfacial oxides MnCr2O4 will decompose, as reported in our previous study [43]. By contrast, the bonding interface formed in the high vacuum of 10-4 Torr is well bonded and no interfacial oxides are observed at the 25
bonding interface. This is because the initial oxide scale on the contacting surface (i.e. “oxidised surface I”) is mainly composed of Fe and Cr oxides and it has a nanoscale thickness. Similarly, once two “oxidised surfaces I” contact with each other, the Fe- and Cr-rich oxide scales are broken up into isolated particles distributed on the bonding
IP T
interface due to deformation. The nanoscale Fe- and Cr-rich oxide particles are unstable and decompose immediately during hot-compression bonding. Thus no interfacial
SC R
oxides exist on the bonding interface without further holding treatments.
U
5. Conclusions
N
The surface chemistry and structure of initial surface oxide scales formed on 316LN
A
stainless steel and the interfacial characteristics after hot-compression bonding under
M
different vacuum conditions were investigated using a combination of XPS, TEM, SEM
ED
and thermodynamic calculations. We report the following findings: (1) The oxide scale of as-received steel consists of binary oxide of Fe and Cr with a
PT
thickness of approximately 4 nm. After holding in a high vacuum of 10-4 Torr, an oxide
CC E
scale with a thickness of 30 nm is observed. While in a low vacuum of 10-1 Torr, the oxide scale is a duplex structure with a thickness of approximately 500 nm.
A
(2) The initial oxide scale formed in a high vacuum of 10-4 Torr is composed of outer
Fe3O4 and inner Fe- and Cr-rich layer. After hot-compression bonding, the bonding interface is well bonded and no interfacial oxides are observed on the interface, owing to the nanoscale thickness and instability of the initial oxide scale. (3) The initial oxide scale formed in a low vacuum of 10-1 Torr is composed of outer 26
MnFe2O4 layer and inner layer consisting of mixed FeCr2O4 and MnFe2O4 along with small amounts of Si-rich nodules. Nickel and molybdenum enrichments are observed at the oxide scale/matrix interface. The thick complex oxide scale on the contacting surfaces transforms into MnCr2O4 after hot-compression bonding, which is resulted
IP T
from the decrease of oxygen activity at the bonding interface. (4) Different vacuum conditions result in different thicknesses and types of oxide
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scales. A higher vacuum leads to a thinner and unstable oxide scale, which decomposes
U
easily during the hot-compression bonding process, leading to a well-bonded interface.
N
Acknowledgements
A
The authors acknowledge the financial support from National key Research and
M
Development program (Grant No. 2016YFB0300401), National Natural Science
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Foundation of China (Grant Nos. U1508215, 51774265) and key Program of the
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