電鍍鋅鎳/鉻複合鍍層於直接式熱沖壓鋼板高溫保護性質研究

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國立臺灣大學工學院材料科學與工程學研究所碩士論文

Department of Materials Science and Engineering College of Engineering

National Taiwan University Master Thesis

電鍍鋅鎳/鉻複合鍍層於直接式熱沖壓鋼板高溫保護性質研究

Electrodeposited Zinc-Nickel/Chromium Binary Coating and Its High Temperature Protection Properties for Direct Hot Stamping

Steels

吳駿泓 Chun-Hung Wu

指導教授：林招松博士 Advisor: Chao-Sung Lin, Ph.D.

中華民國 103 年 7 月

July, 2014

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誌謝

還記得推甄上台大材料所的時候，心中對於要留在清大材料還是匇上來台匇念書仍然猶豫不決。大學就離開家住台南的我，要適應一個新的環境並不會太困難，真正困擾我的是如何做一個不讓自己後悔的決定。大三那年到瑞典林雪帄大學當交換學生的生活讓我更願意在新的事物上做嘗試，因此我毅然決然地匇上到台大來開啟我的研究所生活。還記得初次見到老師時老師非常熱心地向我介紹實驗室在做的研究，讓我對實驗室有個初步的認識。當時找指導教授還發生一個小插曲，當時找完老師正準備離去前，老師看我特地從新竹跑來台匇，問我還有沒

有找其他的教授，我跟老師說：「沒有，我想說若是老師確定不收我後我再去詢問

其他的老師」。結果老師跟我說不確定能不能收我，因為台大材料所找指導教授是

要選填志願序的，而老師當時已經收了一個推甄上的專題生。當下聽到其實蠻沮喪的，本來想說不要抱著走馬看花的心情找指導教授，卻沒想到自己可能因此沒有適合自己的教授能夠收我。非常感謝當時我的女友怡安在一旁聽我分享與安慰鼓勵我並且不斷地為這件事來禱告，最後也如願進到了老師的實驗室當中！

我真的很慶幸自己成為表面處理實驗室當中的一員，第一次見到大家時就發現實驗室的氣氛非常的好，學長姐與同學也都會主動地與我聊天。兩年的時間匆匆地過去，第一次拿到實驗室鑰匙的那份感動仍存在心中。兩年下來累積滿滿的回憶實在難以用紙筆來傳達，不論是抒壓解放的唱歌團、熬夜爆肝的 RM 團、別出心裁的聖誕趴、大吃大喝的實驗室出遊、葡萄美酒的夜光杯以及揮灑汗水的夜光盃，每次的活動都將大家更緊密的連結在一起。這兩年來受過實驗室的各位許多的幫助，謝謝尉桓學長大談軍中的生活點滴，讓我在當兵之前能夠先打一劑強心針；謝謝黼澤學長總是熱心幫助大家，學長的 TEM 技術實在是無人能及；謝謝喻仁學長常常不厭其煩地回答我電化學相關問題，讓我搞懂了很多以前錯誤的觀念；

謝謝憲中學長偶爾會跑回實驗室與我們哈拉，跟我們分享在某公司被凹到爆的生活；謝謝香孙學姊把關實驗室環境整潔，沒有學姊的工綜館在與舊物館相比高下立判；謝謝克駿學長為實驗室帶來歡笑與陽光，讓實驗室時時充滿著溫暖的氛圍；

謝謝順億學長熱心地幫忙打 TEM，還會分享一些股市投資的資訊；謝謝宛珊學姊在每次報告結束後都會再給我一些建議，讓我知道還有甚麼需要加強的地方；謝謝

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謝謝高峰學長分享以前在清大的趣事；謝謝文昕學姊替實驗室節省荷包；謝謝秀瑜學姊教我做實驗；謝謝小巴晚上在實驗室裡一起邊做實驗邊吃宵夜，讓宵夜多了股特別的味道；謝謝茂峰碩二一整年來陪我在工綜度過一年的時光，不論是休閒的韓綜、棒球或是嚴肅的實驗問題及未來規劃都成為我們聊天的話題；謝謝俊銘成為實驗室強而有力的支柱，無論大大小小的事都難不倒你，只可惜你沒有要留下來讓實驗室頓失依靠；謝謝碩一的學弟們：總務大臣漢邦、卷哥代表崧貿、

東區新貴建豪、卡神老闆恆佳、肌肉猛男清華，感謝你們在口試的期間幫了許多忙，包括訂購茶點、場地布置以及記錄問題，祝福你們未來一年內可以在實驗室中有更多的學習與收穫；謝謝工綜的老鼠(們)總是溫馨地提醒我們該回家了，雖然你們有時對待食物的方式不太友善。

在這邊也要好好感謝我的指導老師，雖然老師當系主任後比較忙碌，比較少機會能與老師討論自己的實驗，但老師在我報告時仍會給予我很多的建議，對於我未來打算繼續出國深造一事也給予肯定和支持！另外也要感謝我的家人和朋友，

在我實驗進行得不順利或面臨一些抉擇時陪我一起度過，先謝謝爸媽在我這兩年來台匇念書期間仍不斷關心我的近況，讓我感受到家庭的溫暖；也謝謝我那有美工天分的老弟有時會幫我製作一些圖表，其他則由於要感謝的人太多就不一一列舉了。

最後，我要向我的女朋友怡安獻上萬分的感激，因為有妳我才有努力奮鬥的目標。研究所的日子一路走來有歡笑也有淚水，難過時有妳傾聽心中滿腹的苦水，

開心時與妳分享心中無比的喜悅。很高興我們已經一起完成了碩士的這個階段，

希望未來的日子也能和妳一起攜手向前邁進！

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中文摘要

電鍍鎘保護層在嚴苛海洋腐蝕環境底下具有良好的犧牲保護效果，加上其光亮的金屬光澤外觀，因此廣泛應用於船舶、航太工業的抗蝕鍍層。許多研究發現鎘會對於環境造成危害且有一定的致癌性，近年來則多以鋅或鋅合金鍍層取代之。

鉻為不銹鋼(鉻含量大於 10.5wt%)中主要添加的合金元素，在一般大氣環境下，鉻

可以在表面形成緻密的 Cr2O3氧化層，阻礙鐵離子經由氧化層擴散到表面與氧氣接

觸，進而降低氧化的速率。在本研究當中，使用鋅鎳合金與鉻的雙電鍍層來解決在高強度鋼熱沖壓製程中傳統熱浸鍍鋅層與電鍍鋅層會遇到的液態金屬誘發脆化 (LMIE)與高溫氧化的現象。液態金屬誘發脆化現象(LIME)為鋅鍍層在高溫下熔融成液態，並沿著鐵的晶界擴散進入，導致在沖壓時成為許多應力集中的小區域而造成沿晶破壞；高溫氧化則因鋅為化學活性高的金屬，高溫下會與氧反應形成大量的氧化鋅而使得有效的鍍層量減少並導致組成分布不均。利用掃描式電子顯微鏡(SEM)與穿透式電子顯微鏡(TEM)觀察鋅鎳與鋅鎳/鉻鍍層微結構，並藉由電子微探儀(EPMA)來分析鍍層與鋼底材界面間擴散的情況。高溫拉伸試驗模擬熱沖壓製程中試片在高溫下實際受到張應力的情形，由應力應變圖可以看出不同鍍層對於鋼底材的機械性質有不同的影響。高溫於表面產生的氧化物則利用 X 射線繞(XRD) 射與 X 射線光電子能譜儀(XPS)進行分析，分析的結果可以更了解鍍層在高溫下氧化的情形。綜合上述結果及透過開路電位測試檢視鍍層是否具犧牲保護的效果，

以及評估雙電鍍層(Zn-Ni/Cr)是否適用於實際熱沖壓製程。

實驗結果發現，不同的電鍍溫度與溶液中鎳離子的比例會影響鍍層形貌、組成成分與陰極電流效率，且正常共電鍍的情形多發生於高鎳離子比與高溫(60℃以上)的情形下。EPMA 的結果顯示鋅鎳鍍層在經過熱處理後會造成組成成分改變，

導致此現象的原因與鍍層氧化以及鍍層與底材間的交互擴散有關。將鉻層電鍍於鋅鎳鍍層上方再經熱處理後發現鋅鎳鍍層的組成成分改變現象有明顯的降低。此外，鋅鎳合金層因其高熔點的性質，與純鋅層相比在高溫下與鐵基材的交互擴散較弱，具良好的高溫穩定性。鍍層中鎳的含量上升可提高鍍層的熱穩定性，電化學開路電位量測結果則顯示經熱處理後的鋅鎳鍍層中隨著鎳含量增加其陰極保護的效果隨之下降。在高溫拉伸試驗中，發現在較高的溫度與較慢的拉伸速率下，

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鋅鎳合金鍍層對鋼板的機械性質劣化較純鋅鍍層來得輕微。鋅在鉻中有相當大的

固溶度，在高溫下會擴散至表面形成 ZnO，而 ZnO 與 Cr₂O3在高溫下形成 ZnCr₂O4

的尖晶石結構，而 ZnCr2O4的結構也經由 XRD 與 XPS 鑑別出來，微結構也經由

TEM 進行細部觀察，結果顯示此結構能防止鋅鎳鍍層在高溫下劇烈氧化。從上述結果看來，鋅合金鍍層相較於電鍍純鋅層表現出較高的熱穩定性與較低的腐蝕速率，而鋅鎳/鉻複合鍍層則提供鋅合金鍍層所缺乏的抗氧化能力，可以減少鍍層因高溫下大量氧化所造成的損失。藉由調整適當的鋅鎳鍍層中的鎳含量來避免液態金屬誘發脆化並使鍍層提供足夠的陰極保護能力，鋅鎳/鉻複合鍍層可能會是適合熱沖壓製程的一個保護層。

關鍵字：熱沖壓鋼板、液態金屬誘發脆化、抗高溫氧化、陰極保護、電鍍複合層

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ABSTRACT

Electrodeposited cadmium layer had provided excellent cathodic protection under serious marine conditions and expressed bright metallic color. Nowadays, the electrodeposited zinc and zinc-based coatings have been extensively applied in general auto industries as replacements for cadmium, which was found carcinogenic in many previous literatures. Chromium is the main alloy element to be added in stainless steel to form a compact chromium oxide above the surface, which retards Fe²⁺ ions diffusing through the oxide, thus slows down the oxidation rate. In this work, the electroplated bilayer (Zn-Ni/Cr) was used to prevent the Liquid Metal Induced Embrittlement (LMIE) phenomenon and serious oxidation that typically occur in hot stamping process degrading the mechanical properties and consuming the available coatings. The microstructure of Zn-Ni and Zn-Ni/Cr coatings was observed by SEM and TEM, and the thermal stability was examined by EPMA and Line scan. The mechanical properties of steels with and without zinc-based coating were illustrated by the strain-stress curve of high temperature tensile test and the broken specimens are examined by Metallographic analysis. Moreover, the surface oxides are characterized by X-ray Diffraction (XRD) and X-ray photoelectron spectroscopy (XPS), which help to further understand the oxidation situation in high temperature. By combining the results mentioned above with open circuit test and dynamic polarization curves, it is able to evaluate the performance of Zn-Ni/Cr bilayer coating in simulated hot stamping process.

The results show that Zn-Ni coating’s morphology and current efficiency both strongly depend on the ion ratio in the bath and bath temperature. The normal co-deposition of Zn-Ni always occurs at high nickel ion ratio in the electrolyte and high

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deposition temperature. EPMA mapping results revealed that there’s a serious separation of compositions in the Zn-Ni coating after heat treatment due to the effect of oxidation and inter-diffusion with the substrate. With the addition of chromium layer on top the situation of separation of compositions in the Zn-Ni coating was apparently reduced. The thermal stability is becoming higher when nickel content increases in the Zn-Ni coating; however, the overall reduction potential will also move close to the substrate comparatively. In addition, spinel structure (ZnCr2O4) formed by ZnO and Cr₂O₃ at high temperatures was identified by XRD and XPS measurements, which indicated zinc diffused through the chromium layer to form zinc oxide at the surface.

The phenomenon can be explained by the great solubility of zinc in chromium and high oxygen affinity of zinc. ZnCr2O4 layer was also observed in TEM for detailed information. In conclusion, the Zn-based coatings exhibit higher thermal stability and lower corrosion rate the electro-galvanized coating. On the other hand, Zn-Ni/Cr bilayer provides better oxidation resistance than Zn-Ni coating that reduces the loss of coating through serious oxidation at high temperatures. With the adjustment of suitable nickel contents in the Zn-Ni coating and achieve equilibrium between the prevention of LMIE and enough ability of cathodic protection, Zn-Ni/Cr bilayer might be a solution to meet the requirements of hot stamping process.

Key word: Hot stamping steel; Liquid metal induced embrittlement; High temperature oxidation resistance; Cathodic protection; Electrodeposited multi-layer

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LIST OF FIGURES

Fig.2.1.1 Basic hot stamping process chains: (a) direct hot stamping, (b) indirect hot

stamping ... 4

Fig.2.1.2 (a) Mechanical properties of 22MnB5 (b) CCT diagram……….. ... 4

Fig.2.1.3 Comparison of the different coating systems for hot stamping application .... 10

Fig.2.3.1 Scheme of the reduction reactions that predominate in different potential domains ... 19

Fig.2.4.1 Diagram of separate two reactants ... 20

Fig.2.4.2 (a) p-type(cation mobile) (b) n-type(anion mobile) ... 21

Fig.2.4.3 (a) Metal-excess n-type semiconductor ZnO with Interstitial cations and excess electrons (b) Metal-deficit p-type semiconductor NiO with cation vacancies and positive holes ... 21

Fig. 2.4.4 Simple model for diffusion-controlled oxidation ... 23

Fig. 3.1.1 Experiment procedure ... 26

Fig. 3.2.1 Scheme of pulse current parameters ... 29

Fig. 3.3.1 Thermal/Mechanical Simulator (Gleeble 3500) ... 33

Fig. 3.3.2 Specimens for high temperature tensile test ... 33

Fig. 3.3.3 Theoretical anodic polarization scan. ... 35

Fig. 4.1.1 Electrodeposition specimens：Steel, Zn-Ni(35℃), Zn-Ni(45℃), Zn-Ni(55℃), Zn (from left to right) ... 36 Fig. 4.1.2 Variation of the nickel content and current efficiency with deposition

temperature： (a)40%Ni in the bath (b) 45%Ni in the bath (c)50%Ni in the

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bath ... 39 Fig. 4.1.3 XRD patterns of zinc-nickel alloy coating with various temperatures：

(a)40%Ni in the bath (b) 45%Ni in the bath (c)50%Ni in the bath ... 41 Fig. 4.1.4 Surface morphology of Zn-Ni coating at various temperatures (40% nickel ions in the bath) : low magnitude (a-1)30℃ (b-1)35℃ (c-1)40℃ (d-1)45℃

(e-1)50℃ (f-1)55℃ (g-1)60℃ high magnitude (a-2)30℃ (b-2)35℃

(c-2)40℃ (d-2)45℃ (e-2)50℃ (f-2)55℃ (g-2)60℃ ... 44 Fig. 4.1.5 Surface morphology of Zn-Ni coating at various temperatures (45% nickel ions in the bath) : low magnitude (a-1)30℃ (b-1)35℃ (c-1)40℃ (d-1)45℃

(c-2)40℃ (d-2)45℃ (e-2)50℃ (f-2)55℃ (g-2)60℃ ... 46 Fig. 4.1.6 Surface morphology of Zn-Ni coating at various temperatures (50% nickel ions in the bath) : low magnitude (a-1)30℃ (b-1)35℃ (c-1)40℃ (d-1)45℃

(c-2)40℃ (d-2)45℃ (e-2)50℃ (f-2)55℃ (g-2)60℃ ... 48 Fig. 4.1.7 Cross-sectional morphology of Zn-Ni coating at various temperatures ... 49 Fig. 4.1.8 Plane morphology of Zn-Ni coating at various temperatures：High magnitude (a-1) 35℃ (b-1) 45℃ (c-1) 55℃ Medium magnitude (a-2) 35℃ (b-2) 45℃

(c-2) 55℃ Low magnitude (a-3) 35℃ (b-3) 45℃ (c-3) 55℃ ... 50 Fig. 4.1.9 Plane morphology of Zn-Ni coating at various temperatures after 850℃ heat treatment for 30 minutes：High magnitude (a-1)35℃ (b-1)45℃ (c-1)55℃

Medium magnitude (a-2)35℃ (b-2)45℃ (c-2)55℃ Low magnitude (a-3)35℃ (b-3)45℃ (c-3)55℃ ... 51

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Fig. 4.1.10 Zn-Ni binary phase diagram ... 53 Fig. 4.1.11 Thermal stability test at 850℃ for 30 minutes：(a) Pure Zn (b) Zn-Ni ... 54 Fig. 4.1.12 Line scans of the Zn-Ni coatings obtained at various temperatures after heat treatment for 30 minutes：(a) 35℃ (b) 45℃ (c) 55℃ ... 55 Fig. 4.1.13 XRD patterns of zinc-nickel alloy coating (thick)：(a) Before 850℃ 30 min heat treatment (b) After 850℃ 30 min heat treatment ... 56 Fig. 4.1.14 Mapping of electrodeposited Zn-Ni coating ... 58 Fig. 4.1.15 Macro appearance of specimens after 850℃ 5 min heat treatment：(a) Zn (b) Zn-Ni(35℃) (c) Zn-Ni(45℃) (d) Zn-Ni(55℃) (from left to right) ... 58 Fig. 4.1.16 Mapping of the Zn-Ni coating after 850℃ 5 min heat treatment：(a) 35℃ (b) 45℃ (c) 55℃... 60 Fig. 4.1.17 XRD patterns of the Zn-Ni coating after 850℃ 5 min heat treatment ... 60 Fig. 4.1.18 Back-scattered transverse image of the Zn-Ni coating after heat treatment 61 Fig. 4.1.19 Line scan analysis of the Zn-Ni coating after heat treatment ... 61 Fig. 4.1.20 TEM Cross-section view of Zn-Ni coating after heat treatment ... 63 Fig. 4.1.21 Selected area diffraction (SAD) of Zn-Ni coating after heat treatment ... 63 Fig. 4.1.22 XPS depth profile of the Zn-Ni coatings after 850℃ 5 min heat treatment . 65 Fig. 4.1.23 XPS peak fitting of the Zn-Ni coatings after heat treatment (Zn_2p) ... 66 Fig. 4.1.24 XPS peak fitting of the Zn-Ni coatings after heat treatment (Ni2p) ... 67 Fig. 4.2.1 Cross-section and plane morphology of Cr coating (at fix deposition temperature) on Zn-Ni coating (at various deposition temperatures) ： Cross-section view (a-1) 35℃(Zn-Ni) (b-1) 45℃(Zn-Ni) (c-1) 55℃(Zn-Ni) Low magnitude (a-2) 35℃ (Zn-Ni) (b-2) 45℃(Zn-Ni) (c-2) 55℃(Zn-Ni) High magnitude (a-3) 35℃(Zn-Ni) (b-3) 45℃(Zn-Ni) (c-3) 55℃(Zn-Ni) . 69

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Fig. 4.2.2 Cross-section and plane morphology of Cr coating (at fix deposition temperature) on Zn-Ni coating (at various deposition temperatures) after 850℃ 5 min heat treatment：Cross-section view (a-1) 35℃(Zn-Ni) (b-1) 45℃(Zn-Ni) (c-1) 55℃(Zn-Ni) Low magnitude (a-2) 35℃ (Zn-Ni) (b-2) 45℃(Zn-Ni) (c-2) 55℃(Zn-Ni) High magnitude (a-3) 35℃(Zn-Ni) (b-3) 45℃(Zn-Ni) (c-3) 55℃(Zn-Ni) ... 70 Fig. 4.2.3 Cross-section and plane morphology of Cr coating (at fix deposition temperature) on Zn-Ni coating (at various deposition temperatures) after 850℃ 10 min heat treatment：Low magnitude (a-1) 35℃(Zn-Ni) (b-1) 45℃(Zn-Ni) (c-1) 55℃(Zn-Ni) Medium magnitude (a-2) 35℃ (Zn-Ni) (b-2) 45℃(Zn-Ni) (c-2) 55℃(Zn-Ni) High magnitude (a-3) 35℃(Zn-Ni) (b-3) 45℃(Zn-Ni) (c-3) 55℃(Zn-Ni) ... 71 Fig. 4.2.4 XRD patterns of Zn-Ni/Cr binary coating ... 72 Fig. 4.2.5 XRD patterns of Zn-Ni/Cr binary coating after 850℃ 5 min heat treatment. 72 Fig. 4.2.6 Plane morphology of Zn-Ni/Cr bilayer with cracks in Cr layer with various heat treatment times at 850℃：(a-1) After 5 min heat treatment (low magnitude) (b-1) After 5 min heat treatment (high magnitude) (a-2) After 10 min heat treatment (low magnitude) (b-2) After 10 min heat treatment (high magnitude) ... 73 Fig. 4.2.7 Mapping of the Zn-Ni/Cr binary coating ... 75 Fig. 4.2.8 Macro appearance of specimens after 850℃ 5 min heat treatment：

(a) Zn-Ni(35℃)/Cr (b) Zn-Ni(45℃)/Cr (c) Zn-Ni(55℃)/Cr (from left to right) ... 75 Fig. 4.2.9 EPMA element mapping of the Zn-Ni/Cr binary coatings after 850℃ 5 min

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heat treatment： (a) 35℃ (b) 45℃ (c) 55℃... 77 Fig. 4.2.10 Back-scattered transverse image of the Zn-Ni/Cr coating after heat treatment ... 78 Fig. 4.2.11 Line scan analysis of the Zn-Ni/Cr coating after heat treatment ... 78 Fig. 4.2.12 TEM cross-section view of Zn-Ni/Cr binary coating after heat treatment ... 80 Fig. 4.2.13 Selected area diffraction (SAD) of Zn-Ni/Cr coating after heat treatment .. 81 Fig. 4.2.14 XPS depth profile of the Zn-Ni/Cr coatings after 850℃ 5 min heat treatment ... 83 Fig. 4.2.15 XPS full survey of the Zn-Ni/Cr coatings after 850℃ 5 min heat treatment：

Zn-Ni/35℃(b) Zn-Ni/45℃ (c) Zn-Ni/45℃ ... 85 Fig. 4.2.16 The variation of total thickness before and after 850℃ 5min heat treatment (HT) with Zn-Ni and Zn-Ni/Cr specimens (Calculated by EPMA Line Scan).

... 86 Fig. 4.2.17 The variation of nickel contents in the remained zinc-nickel coating with and without chromium layer after 850℃ 5min heat treatment (HT) ... 87 Fig. 4.3.1 High temperature tensile test at different temperatures and strain rates ... 89 Fig. 4.3.2 OM photos of broken specimens obtained after high temperature tensile test90 Fig. 4.3.3 EPMA mapping of the broken specimens after high temperature tensile test：

(a) Zn：850℃, 0.26s^-1 (b) ZnNi：850℃, 0.26s^-1 ... 92 Fig. 4.3.4 EPMA mapping of the broken specimens after high temperature tensile test：

(a) Intragranular fracture specimen (b) Intergranular fracture specimen ... 93 Fig. 4.4.1 Open circuit potential (OCP) of the Zn-Ni/Cr binary coating (V.S SCE)：

(a) before 850℃ 5 min heat treatment (b) after 850℃ 5 min heat treatment95 Fig. 4.4.2 Potentiodynamic polarization curves of Zn-Ni coatings with various nickel

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contents. Scan rate：0.5 mV s^-1 ... 97

Fig. 5.3.1 Cr-Zn binary phase diagram ... 106

Fig. 5.3.2 Ni-Cr binary phase diagram... 107

Fig. 5.3.3 Crystal structure of spinel ZnCr₂O₄ ... 107

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LIST OF TABLES

Table.2.1.1Chemical components and mechanical properties of boron steels ... 5

Table.2.4.1 Pilling-Bedworth ratios R for various metal/metal oxides ... 25

Table 3.1.1 Components of commercial low carbon steel ... 26

Table 3.2.1 Composition of Zn-Ni electrolyte ... 27

Table 3.2.2 Composition of Cr(Ⅲ) electrolyte ... 28

Table 4.1.1 Composition of Zn-Ni coating after 850℃ 5 min heat treatment ... 62

Table 4.2.1 Composition of Zn-Ni/Cr coating after 850℃ 5 min heat treatment ... 79

Table 4.4.1 Electrochemical corrosion parameters of different specimens ... 97

Table 5.3.1 Distribution of electrons in 3d orbitals and CFSE of first transition series ions in octahedral coordination (high spin state) ... 108

Table 5.3.2 Crystal field stabilization energies (kJ/mole) of transition metal ions ... 108

Table 5.5.1 Standard Reduction Potential (obtained at 298.15K and 101.325kPa)... 111

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Chapter 1 Introduction

In order to achieve the demands of vehicle weight reduction, passenger’s safety improvement and crashworthiness qualities, the need for manufacture automobile structural components from advanced-high-strength steels (AHSS) has increased. Cold forming process for AHSS such as dual-phase steel (DP) and twinning-induced plasticity steel (TWIP) are sensitive to springback damage. Instead, hot stamping process provides excellent formability and no springback damage. On the other hand, the production of hot-stamped car parts for body in white has drastically increased in recent years. Moreover, in comparison to conventional cold forming, there are advantages of better accuracy of size and less forming pass in hot stamping process.

Many studies on hot-dip galvanizing zinc coating for hot stamping sheet steel were investigated, and the main problems faced are liquid-metal-induced embrittlement (LMIE) in the early stage of the press forming at high temperature and high temperature oxidation. LMIE was due to liquid phase penetration in the grain boundary of steel substrate, and high temperature oxidation was caused by high oxygen affinity of zinc under high temperatures. In hot stamping process, the steel blank is heated in a furnace for a few minutes (3-10min) in a temperature range of 850-950℃ in order to get fully austenitic steel, then it is quickly transferred to a press forming a desired component shape and simultaneously hardened by die quenching (cooling rate higher than 50℃/s to avoid the bainite transformation in the deformed areas). Due to high heating rate and high processing temperature, zinc (melting point 419.5℃) will inevitably melt and then penetrate into the grain boundary of iron by fast inter-diffusion at high temperatures.

This phenomenon would cause induced crack propagation through grain boundaries during the press, the bare steel substrate is then exposed to corrosive environment and

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loses the cathodic protection of hot-dip galvanizing coating. Serious oxidation and decarburization also occur at high temperatures during the pre-heat treatment process, which will both reduce the thickness of coatings and the strength near the surface of substrate.

Although hot-dip aluminized coatings are good solutions to solve LMIE phenomenon and serious oxidation at high temperatures in many literatures; however, the coatings do not provide cathodic protection, only high barrier protection. To achieve adequate corrosion protection compared to galvanized or galvannealed parts produced by means of cold forming, steel producers are working on the development of suitable zinc based coating system for hot forming processes. The cathodic protection is essential for the anti-corrosion auto-industry, which extends the operational life of the steel. Cadmium coatings are used to enhance the corrosion resistance of different materials on a wide range of components and parts in many industry applications.

However, Cadmium was found carcinogenic and toxic in many reports and literatures [75-76], so its use has been restricted and replaced by zinc and zinc-based coatings nowadays. In addition, electroplated Zn-Ni or Zn-Co coatings exhibit high thermal stability (high melting point) and are capable to provide cathodic protection, but the high temperature oxidation resistance are relatively poor compared with hot-dip aluminized coatings. In fact, the coating for hot stamping use should simultaneously contain the characters of high thermal stability, great resistance to oxidation and enough cathodic protection to substrate. The present study combines the advantages of high melting point Zn-Ni coating and the formation of a compact oxide of chromium coating, with the simulated thermal, mechanical and electrochemical tests, aiming at evaluating whether the bilayer coating is suitable for hot stamping process or not.

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Chapter 2 Literature Review

2.1 Hot stamping

2.1.1 Introduction to hot stamping

The technique of hot stamping was first invented and patented by a Swedish company (Plannja) to produce saw blades and lawn mower blades [1]. Saab Automobile AB was the first vehicle manufacturer who used a hardened boron steel component in Saab 9000 [2]. There are two types of hot stamping processes: the direct and indirect hot stamping. According to the literature [3], a blank is heated up in a furnace, transferred to the press and subsequently formed and quenched in a specific die in the direct hot stamping process (see Fig. 2.1.1(a)). On the other hand, the indirect hot stamping process is characterized by the use of a nearly complete cold pre-formed part which is subjected only to a quenching and calibration operation in the press after austenitization (see Fig. 2.1.1(b)). The blank exhibits a ferritic–pearlitic microstructure with an initial tensile strength of about 600 MPa in the beginning. After the hot stamping process, the component finally has a martensitic microstructure with a total strength of about 1500MPa as shown in Fig. 2.1.2 (a). To achieve such a high hardness microstructure, take the 22MnB5 steel for instance, the cooling rate must exceed the minimum cooling rate of approximately 27K/s, which leads to induce the diffusionless martensitic transformation at about 400 ℃ (see Fig. 2.1.2 (b)) and produces the parts with high strength [3]. Naderi [5] investigated the mechanical properties of ultra high strength steels produced by hot stamping process and the critical cooling rates for martensitic transformation, Table 2.1.1 shows that only 22MnB5, 27MnCrB5 and 37MnB4 form fully martensitic microstructure after hot stamping when a water-cooling die is used.

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Fig. 2.1.1 Basic hot stamping process chains: (a) direct hot stamping, (b) indirect hot stamping [3].

Fig. 2.1.2 (a) Mechanical properties of 22MnB5 (b) CCT diagram [4].

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Table. 2.1.1 Chemical components and mechanical properties of boron steels [5].

2.1.2 Liquid Metal Induced Embrittlement (LIME)

Zinc-based coatings for hot stamping steel sheets were developed, which provide additional benefit of cathodic protection. During hot stamping process, for the high temperatures (850℃~950℃) reached at the steel sheet’s surface, the zinc coating applied for corrosion protection would inevitably melt due to its relatively low melting point (~420℃). The presence of liquid zinc combined with the presence of high stresses generated by the thermomechanical welding cycle could result in the liquid metal induced embrittlement (LMIE) phenomenon, which cause a catastrophic deterioration of the material’s mechanical properties [8, 22-23]. Hence, many attempts were made to avoid LMIE including indirect hot stamping, increasing the holding time during austenization, the use of high melting point coating and decreasing hot stamping temperature. In the indirect hot stamping method, there is an additional cold preforming step at room temperature prior to austenization [7]. Since most of the deformation completed before heating, LMIE is not observed during the indirect hot stamping.

However, indirect hot stamping is not suitable for continuous galvanizing process,

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usually for batch galvanizing process. Moreover, there will be an extra cost of the mold components for indirect hot stamping. On the other hand, Lee et al. [8] reported that increasing austenization times was also a suitable method to avoid Zn-induced LMIE.

Through the interaction between steel and galvanized coating at high temperatures, liquid zinc or liquid Fe-Zn intermetallic phase will fully transform into α-Fe(Zn) phase with the increment of dwell time. However, the increasing dwell time also leads to serious oxidation that consumes zincs of galvanized coating. Kondratiuk et al. [9]

indicated the addition of 11 wt% Ni increases the melting temperature of Zn-Ni intermetallics (880℃), which is higher enough to prevent LMIE. Nevertheless, The scale formed on the Zn-Ni surface cannot retard the oxidation kinetics as an alumina layer. Abbasi et al. [10] and Fan et al. [11] pointed out LMIE can be avoided due to the absence of a liquid phase, through a preious deformation below the melting temperature of the Fe-Zn intermetallics, but this approach leads to the transformation from the austenite to strain-induced ferrite and bainite, and it will result in a considerable reduction of the strength of the hot stamping parts. Gu et al. [12] studied the structure and phase transformations of electroplated zinc and zinc–iron deposits after heat treatments. The results revealed that the plane morphologies of zinc and zinc-iron coatings did not change with temperatures until reaching a critical point. The critical temperature increases with more iron contents in the coating. Once heating above the critical point, the surface becomes ruffled. In Bories’ study [13], it was shown that the thermal stability of Zn-Ni electrodepsited layer containing less than 13wt% Ni is insufficient, and will result in the lack of corrosion protection in thermally selected bodies. However, the zinc based alloy coating systems for hot forming applications only limited literature has been published [9, 77-78], especially on coating behavior above 700 °C.

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2.1.3 Surface oxidation and decarburization

Other serious problems encountered during hot stamping are surface oxidation and decarburization. Decarburization becomes serious when the metal is heated to temperatures of 700°C or above when carbon in the metal reacts with gases containing oxygen or hydrogen [14]. Imai et al [15] investigated galvannealed steel sheets heated to 900 °C for 5 minutes and found that the coating completely transformed into a zinc oxide layer and there’s iron zinc solid solution containing 20–30 wt.% zinc left after hot stamping while iron oxide was not observed on the surface. Therefore, most sheet metal blanks are pre-coated with a protective layer to solve these problems. Al–Si coating layer is commonly used for preventing scale formation on the steel during the direct hot stamping operation [16-20]. Borsetto et al. [20] studied the influence of thermal process parameters on the chemical behavior of the coating of the Al–Si layer, and the results showed that this protective layer prevents the formation of scales in the direct hot stamping process. Lee also reported that thin Al-10 pct Si alloy (~25) coating help to prevent the steel from decarburization and surface oxidation during the high-temperature thermal cycle in hot stamping process. However, due to the lower forming limits of the Al–Si layer compared to the base material in the initial state at room temperature, the hot-dip aluminized sheets cannot be used for the indirect process and they are not suitable for cold forming. Moreover, Al-Si coating does not provide cathodic protection, like zinc or cadmium, but possesses a high barrier protection. From another aspect, according to Mori and Ito [21], two kinds of different oils were applied for the prevention of oxidation, which were conducted and evaluated in a cooling experiment without forming and in a hot bending experiment. The examination of the sheet surface showed that the number of lubrications (up to 4 times) reduced the

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either adding suitable elements that will form compact scales (such as alumina and silica) above in the coating or covering a protective barrier layer (ex. Painting or Oils) will help to reduce the level of oxidation.

2.1.4 Different types of protective coatings for hot stamping process

De Cooman et al. [78] reviewed different coatings for hot stamping use, including Aluminized coating, Galvanized coating (GI), Galvannealed coating (GA), Zn-Ni coating and Hybrid coating (see Fig. 2.1.3). It can be noticed that the research of Zn-Ni as a protective coating for hot stamping published are comparatively fewer than other coatings. The aluminized coating forms two different structures through different heat treatment methods, type 1 exhibits layered structure containing Fe₂AlSi₂ phase and Fe2Al5 phase while type 2 contains ductile Fe3Al phase and FeAl phase. In addition, the Fe-Zn intermetallic compounds and α-Fe(Zn) phase can be obtained by galvanized and galvannealed coating through the holding time during heat treatment. The hot stamping related characteristics of protective coatings are listed in Table 2.1.2 (The hybrid coating is not discussed in this study). Both two types of aluminized coatings exhibit excellent oxidation resistance by forming compact and thin Al2O3 layer on the surface. GI and GA are also reported to have enough oxidation resistance due to the trace of Al in the coating diffuses to the surface and form Al2O3. Therefore, the oxides on the GI and GA surface after hot stamping contain Al₂O₃ and ZnO. On the other hand, the Zn-Ni coating exhibits the poorest oxidation resistance among the protective coatings. No Al in the Zn-Ni coating results in large amounts of ZnO on the surface after hot stamping. Liquid metal induced embrittlement (LMIE) is commonly observed in two metals with low mutual solubility. The solubility of Zn in α-Fe is high while it is relatively low in γ-Fe.

When the temperature is above peritectic temperature, there will be liquid zinc present.

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If liquid zinc contact withγ-Fe, it will lead to the embrittlement of austenitic steel by facilitating the decohesion of grain boundaries [79]. There’s no LMIE found in aluminized coatings for the rapid formation of high melting point Fe-Al intermetallic compounds. The Zn-Ni coating can avoid LMIE by increasing nickel content in the coating. LMIE phenomenon is commonly seen in GI and GA coatings where liquid zinc might present during hot stamping. The use of indirect hot stamping process or increasing the holding time at high temperatures are developed to prevent LMIE on GI and GA specimens [8, 22-23,78, 80]. Aluminized coatings provide barrier protection by forming stable corrosion products while GI, GA and Zn-Ni coatings provide cathodic protection due to its more negative electrochemical potential than steel. Since the boiling point of Zn (907°C) is relatively low, the coating evaporation will become serious when heating above the boiling temperature of Zn. Therefore, the evaporation problem occurs in GI and GA coatings while in Zn-Ni coating is suppressed by stronger Zn-Ni bonding. This problem is not found in aluminized coatings due to the higher boiling point of Al (2519°C). The layered structure of type 1 aluminized coating is good for resistance spot welding in Drillet’s study [81], which is reported to have a wide welding current range of 1.4 kA. Faderl [80] studied the RSW properties of GI specimen after hot stamping. It was found that the surface oxides on the surface will increase the contact resistance of RSW. This result was also found in Genderen’s study on GA stamped specimens [82]. Therefore, the surface oxides must be eliminated from the coating prior to welding. The Zn-Ni also need to clean the surface oxides to reduce the contact resistance, but it was reported that Zn-Ni coating has lower electrode wear and wider welding current range than that of GI and GA coatings. In the aspect of paintability, it is necessary for GI and GA to be phosphated before painting [80] while aluminized coating can be painted directly without phosphate layer [83]. On the other

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hand, the formability of type 2 aluminized is better than type 1 for the presence of ductile Fe3Al and FeAl phases. Both Fe-Zn and Zn-Ni intermetallic phases are brittle at room temperature. The α-Fe(Zn) phase is ductile under tensile stress at high temperatures, but the study of Fe-Zn-Ni γ phase are not reported in papers.

Fig. 2.1.3 Comparison of the different coating systems for hot stamping application [78]

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2.2 Anomalous co-deposition theory

According to Brenner’s definition proposed in 1963 [24], since the less noble metal zinc deposits preferentially on the cathode and its percentage in deposit is higher than that in the electrolyte, which typically occurs in zinc-iron group metal co-deposition.

Based on electrochemical theory, the nickel deposition should prior to zinc in the co-deposition of zinc–nickel alloy since the equilibrium potential of nickel is far more positive than zinc; however, many research found that zinc deposits preferentially in most practices [25-28]. Although there are many attempts made to explain the anomalous co-deposition, still no universal theory accepted due to the complicated kinetic process.

2.2.1 Hydroxide suppression theory

One model proposed that anomalous co-deposition was attributed to the increase of Ph value at the cathode surface being able to induce zinc hydroxide precipitation, which inhibits nickel discharge [29-32]. In Dahms’ research [31], the experiment was conducting with dropping mercury as cathode. The result showed that the anomalous co-deposition was observed at high overpotential deposition region, which might due to the hydrogen evolution of proton causing the increase of Ph value, thus zinc hydroxide formed on the cathode. In order to prove this hypothesis, Dahms also did the experiment of single iron group metal deposition, and the results showed that deposition rate was not hindered by its hydroxide. As the critical Ph for precipitation of iron-group metal hydroxides is significantly higher than for precipitation of zinc hydroxide, the former may not form so that M-deposition requires direct discharge of M²⁺ ions through the zinc hydroxide film. Fukushima’s group [30, 33-35] also studied the mechanism of

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anomalous co-deposition, and they found two strong evidences to support hydroxide suppression model. First, the Ph value on the cathode increase drastically when the current density achieves the critical value, which causes the anomalous co-deposition.

Second, there’s a sharp increase of the impedance between cathode and reference electrode. However, there are still some experimental results contrary to this theory. In Fabri Miranda’s study, the zinc content in coatings decreased with the rise of Ph value of zinc-nickel alloy electrolyte. In addition, this theory does not explain the strong inhibition of nickel reduction observed in the normal deposition region, the high current efficiency during anomalous codeposition and the increase in nickel content in the alloy with increasing pH value.

2.2.2 Underpotential deposition (UPD) theory

Another theory assumes that underpotential deposition (UPD) of Zn provides an alloy surface that is different from the parent metal for the continuous codeposition, which means thepreferential deposition of zinc is due to the underpotential deposition of zinc on the surface of iron group metals since zinc deposition potential on surfaces of iron group metals is much higher than that on surfaces of the other metal, thus making the deposition of the less noble component preferable [36-37]. However, if the model was correct, once a monolayer is deposited, the UPD should be terminated and the ions in solution should “sense” only the last layer deposited on the surface. Hence, such a theory is valid only if an alternating multilayer coating is formed, but it cannot explain how zinc deposits continuously when the surface of zinc–nickel coatings is entirely occupied by zinc atoms. Moreover, this theory also cannot describe the preferential deposition of nickel in some specific conditions, such as low current density and low deposition potential [29, 38-40].

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2.2.3 Exchange current density theory

According to a third theory, the great difference between the exchange current densities of Zn and the iron-group metal results is a significant difference between the thermodynamic and the practical nobility. Therefore, the magnitude of the exchange current density is generally much greater for Zn compared to Ni, Co and Fe [36, 41-43].

In Hegde group’s study [44], they found this model more appropriate to explain the deposition behavior in the Zn–Ni-Co system. It should be noted that such a model may be proper for electroplating under galvanostatic conditions in their study, where a high current consumption by one element must be at the expense of another element, but their results may not be applicable to electroplating under potentiostatic conditions .

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2.3 Factors of Zn-Ni co-deposition process

Generally speaking, the alloy co-deposition process is affected by many factors, including bath composition, temperature, additives, current density, deposition potential and bath Ph [38, 45–50]. Mathias et al. [51] studied the effect of bath composition, temperature, electrolyte velocity, current density on the alloy composition and current density distribution within the cell. The experimental results showed that composition uniformity could be achieved when the mass transfer was fast relative to the electrode kinetics, so that the surface concentrations of the reacting species remained essentially at their bulk values. Therefore, the temperature and bath composition could affect the average alloy composition through their influence on the electrode kinetics. In conclusion, an acceptable theory would be developed if the zinc–nickel co-deposition mechanism under various deposition conditions is made certain.

2.3.1 Deposition temperature

The effect of temperature on the electrodeposition of zinc–nickel alloy from alkaline bath and acid bath had been studied [39, 52-53]. The structure, composition, mechanical, optical and thermal property of electroplating alloys has strong relationships to deposition temperature. In most cases, the deposition rate is also subject to deposition temperature since the diffusion of metal ion from bulk to cathode will be accelerated with the rise of deposition temperature. With Lee’s group studied Zn-Ni co-deposition in alkaline bath [39], Abou-Krisha’s group in sulphate bath [52] and Qiao’s group [53] in chloride bath, the results of them showed that the nickel content in zinc–nickel alloy coating was dependent on the deposition temperature when other plating parameters were fixed. The nickel content in the coating of three different

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electrolytes went up with the increase of temperature, which might due to the result of intrinsically slow nickel kinetics. According to the literatures [38-40], it is suggested that the preferential deposition of nickel is attribute to a mixed intermediate ZnNi⁺_ads which catalyzes the reduction of nickel ion. The preferential deposition of zinc is attributed to the intermediate Zn⁺_ads which play a catalytic role on the deposition of zinc rich coatings. Qiao’s research group pointed out that the formation of Zn⁺ads on the surface of cathode accelerating the reduction of zinc and blocking the nickel to deposit at low temperature. On the other hand, once the temperature increases, the formation of mix intermediate ZnNi⁺ads dominates the reduction process, which catalyzes nickel deposition and hydrogen evolution on the surface of cathode.

2.3.2 Bath content ratio

Rehim et al. [32] observed that the Ni content in the deposits decreases with increasing Zn content or with decreasing Ni content in the bath when other parameters are fixed, which might due to the fact that an increase in metal content in the bath tends to oppose the depletion of that metal in the cathodic diffusion layer. Roventi et al. [40]

studied the relationship between the deposition potential and bath content ratio. The results showed that at a potential of -700 Mv the alloy composition is almost constant and does not seem to depend on the bath composition; while deposition potential decreases, the nickel contents in the deposits rises with the increase of Ni²⁺ ion content percentage in the bath. Byk et al. [56] investigated the influence of different bath content ratio on compositions in the deposits, deposition rate and current efficiency in chloride bath and ammonical diphosphate bath. The deposition rate and current efficiency reduces with Ni²⁺ ion content percentage increase in the bath, while the

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ammonical diphosphate bath, with the change is sharper in ammonical diphosphate bath than in chloride bath.

2.3.3 Additives

Albalat et al. [45] have examined the relationship between additives and deposit properties in Zn-Ni coatings. In their work, the electroplating of Zn-Ni alloys from a chloride bath containing two brighteners (phenolic derivative and an unsaturated aromatic compound) and a leveling agent (an aromatic carboxylate) has been studied under different plating conditions. According to the results, the composition and morphology of the alloys deposited depended on the concentration of all the additives and also on the temperature. On the other hand, from their results it could be concluded that the corrosion resistance of the coating was related more to their morphology than to their composition. The best behavior could be obtained with alloys that had nodular grains of measurable size, while those that had elongated or non-measurable grains always gave lower corrosion resistance.

2.3.4 Electrolyte type

Many studies of Zn-Ni co-deposition system were conducting in chloride bath [26, 29, 40, 51, 53-54, 56] and in sulfate bath [30, 34, 37, 55]. One reaction model has been proposed by Chassaing and Wiart [54] in chloride bath. A mixed compound (ZnNi⁺ads) dominates the deposition of zinc-rich alloys. This compound plays a role as a catalyst for nickel deposition, and is incorporated in the alloy co-deposition with increasing polarization, thus allowing zinc to deposit preferentially. At low cathodic polarization, which normal co-deposition usually occurs, the deposition of nickel-rich alloys was

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attributed to a mixed intermediate (ZnNi⁺ads), which catalyses the reduction of Ni²⁺ ion.

On the contrary, at high cathodic polarizations, which anomalous co-deposition often occurs, zinc preferential discharge is attributed to the intermediate Zn⁺_ads, catalyst for the deposition of zinc rich deposits. Metal cations can easily form chloro-complexes in chloride electrolyte [51]; however, it should be considered the reactive species in sulfate electrolyte are Ni²⁺ and Zn²⁺ ions. There are four factors found in Fabri Miranda’s research in sulfate bath [55], (1) ZnNi⁺ads acts as a catalyst in nickel deposition (2) The intermediate H_ad involved in hydrogen evolution at low polarizations (3) The intermediate adion Zn⁺ads acts as a catalyst in zinc deposition at high polarizations (4) The anionic Zn/NiSO₄^2-_ads inhibits zinc deposition at intermediate polarizations through the incorporation of S deposition.

Chassaing’s reaction model in chloride bath [54]

Low polarizations and high temperature: normal deposition and hydrogen evolution

2

incl.

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+ 2

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2

H 2H

Zn + (NiZn)

+ 2e + Zn + (NiZn)

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High polarizations and low temperature: anomalous deposition

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Fabri Miranda’s reaction model in sulfate bath [55]

Low polarizations and high temperature: normal deposition and hydrogen evolution

(7.2)

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Transition region

(11) Zn(OH)

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High polarizations and low temperature: anomalous deposition

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Fig. 2.3.1 Scheme of the reduction reactions that predominate in different potential domains [55].

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2.4 High temperature oxidation theory

From simple oxidation reaction:

) ( )

( 2 )

(

2

1

S g

S

O MO

M  

The solid scale MO will separate the two reactants as shown below in Fig. 2.4.1

Fig. 2.4.1 Diagram of separated two reactants (metal and gas) [57].

According to the textbook [57], for the oxidation to move on, at least one reactant need to penetrate the scale, either metal must be transported through the oxide to the oxide–gas interface and react there, or oxygen must be transported to the oxide–metal interface and react there.

There are two types of the scale are shown in Fig.2.4.2, which also presents two oxidation mechanisms. The growth of the scale of these two types is different. Cation migration from substrate leads to scale formation at the scale–gas interface whereas anion migration from atmosphere leads to scale formation at the metal–scale interface.

In most cases, to describe simultaneous migration of ions and electrons it is necessary to assume that the oxides are non-stoichiometric compounds. Non-stoichiometric ionic compounds are classified as semiconductors with the chemical formula M1+δO and M_1−δO (the value of δ varies widely from 0.05 to 0.001). NiO, FeO &Cr₂O₃ are p-type cation-deficit semiconductors; therefore, the cations will migrate with electrons from the scale–metal interface to the scale–gas interface during oxidation. On the other hand,

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ZnO is an n-type cation-excess semiconductor, having interstitial Zn ions and equivalent electrons within the conduction band. The oxidation occurs when anions migrate from the scale–gas interface to the scale–metal interface. The scheme of ZnO and NiO are shown in Fig. 2.4.3.

Fig. 2.4.2 Interfacial reactions and transport processes for high-temperature

oxidation mechanisms (a) p-type(cation mobile) (b) n-type(anion mobile) [57].

(a) (b)

Fig. 2.4.3 (a) Interstitial cations and excess electrons in ZnO – an n-type metal-excess semiconductor. (b) Typical p-type metal-deficit semiconductor NiO with cation vacancies and positive holes. [57].

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When two reactants are separated by the scale produced, it is necessary to postulate that ionic and electronic transport processes through the oxide are accompanied by ionizing phase-boundary reactions and formation of new oxide at a site whose position depends on whether cations or anions are transported through the oxide layer.

Take a simplified treatment of diffusion-controlled oxidation in textbook [57].

Assume all thermodynamic equilibriums are established at each interface, and the cationic transport through the scale governs the rate of oxidation, the process are listed below. The outward cation flux, j_M²⁺, is equal and opposite to the inward flux of cation defects (here taken to be vacancies). This model is shown in Figure 2.4.4.

Thus, j_M²⁺ can be expressed as in Equation: [2.3.1]:

C' - ' D C' j -

j

^M ^M

M + M

2

V V

V M V

 x



[2.4.1]

where x is the oxide thickness, D_VM is the diffusion coefficient for cation vacancies, and C’VM and C’’VM are the vacancy concentrations at the scale–metal and scale–gas interfaces, respectively.

Since there is thermodynamic equilibrium at each interface, the value of (C’’_VM – C’_VM) is constant and we have Equation [2.3.2]:

x

M M

+ M 2

V V

V OX

M

C' - ' D C'

dt dx V

1 j  

[2.4.2]

i.e., the rate of oxide thickening is given in Equation [2.3.3]

) ' ''

( '

'

M M

M OX V V

V

V C C

D k where x

k dt

dx   

[2.4.3]

where Vox is the molar volume of the oxide. Integrating and noting that x = 0 at t = 0 we obtain Equation [2.3.4]:

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x

²

 2k' t

[2.4.4]

which is the common parabolic rate law.

Furthermore, since it has been shown that the cation-vacancy concentration is related to the oxygen partial pressure by Equation [2.3.5]:

1n O VM

const. P

2

C 

[2.4.5]

the variation of the parabolic rate constant with oxygen partial pressure can be predicted, i.e., Equation [2.3.6]:

] ) (p' - ) ' (p' [

'

_O ¹ⁿ _O ¹ⁿ

2



2

k

[2.4.6]

Since p’O2 is usually negligible compared with p’’O2 we have Equation [2.3.7]:

' ( p ' '

_O

)

¹ⁿ



2

k

[2.4.7]

Fig. 2.4.4 Simplified model for diffusion-controlled oxidation [57].

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2.4.1 Pilling-Bedworth ratio

Pilling-Bedworth ratio (PBR) is a suitable tool to evaluate the resistance of a material against high temperature oxidation. [58] The performance of protective oxide barriers usually depend on the continuity and the compactness. The protective layer might be deformed by the generation of cracks, which then results in cracks and spallation. There are two types of stress occurred during oxidation, growth stress and thermal stress. The former is affected by the growth rate of oxide, while the later is caused by the difference in thermal expansion coefficient. PRB can be expressed as below :

Tensile strength will be generated in the oxide if PBR<1, which means metals cannot produce sufficient oxide to cover the surface. On the other hand, compress stress will be generated in the oxide when PBR>1, and there will be favorable compressive stress in the oxide if PBR is slightly higher than 1 while cracks might be formed if the values of PBR is much higher than 1. Generally speaking, the value of PBR between 1~2 exhibits more compact oxide and better resistance to corrosion due to its moderate compressive stress. Table 2.4.1 presents the PBR values of various metal/metal oxides.

It can found that alkali metals and alkaline earths have PBR values smaller than 1, and iron oxides such as Fe₂O₃ and Fe₃O₄ have PBR values higher than 2. Although the effects of PBR on the stress are complicated and not well studied, the application of PBR is more important in high temperature oxidation where thicker oxides exhibit mechanical and physical properties close to bulk oxide.

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Table. 2.4.1 Pilling-Bedworth ratios R for various metal/metal oxides [59].

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Chapter 3 Experimental method

3.1 Experimental design

The experiment procedure is shown in Fig.3.1.1, and the specimens used are commercial low carbon steel, which components are listed in Table.3.1.1.

Fig. 3.1.1 Experiment procedure.

Table 3.1.1 Components of commercial low carbon steel.

C Si Mn P S

Wt% 0.04 0.002 0.18 0.014 0.015

電鍍鋅鎳/鉻複合鍍層於直接式熱沖壓鋼板高溫保護性質研究

國立臺灣大學工學院材料科學與工程學研究所 碩士論文

Department of Materials Science and Engineering College of Engineering

National Taiwan University Master Thesis