鎳基高溫合金在高溫具有優(yōu)異的力學和抗氧化腐蝕性能，是制造航空發(fā)動機和燃氣輪機核心零部件的關(guān)鍵材料[1] 隨著對航空發(fā)動機性能要求的提高，對其承溫能力提出了更高的要求 高溫合金中的難熔元素(W、Mo、Re等)能有效地提高其固溶強化水平，從而使其承溫能力提高[2~5] 用W元素部分替代高溫合金中的Re等稀貴金屬元素，在提高其力學性能的同時還能降低制造成本[6] 因此，對高鎢高溫合金的研發(fā)受到國內(nèi)外學者的極大關(guān)注[7, 8]

鎳基高溫合金中的析出相主要有A1型γ基體和L12型γ?有序相，皆為面心立方結(jié)構(gòu)且有明顯的共格關(guān)系[1] γ?相是鎳基高溫合金中最重要的強化相，其含量、形貌和尺寸都顯著影響合金的高溫力學性能[9, 10] 因此，在高溫合金的變形過程中位錯與γ?相的交互作用極為重要 Wang等[11] 的研究表明，在室溫拉伸過程中位錯以層錯的形式切入γ?相并阻礙其他位錯運動，使合金的屈服強度提高 Cui等[12]發(fā)現(xiàn)，M951G合金中的連續(xù)層錯穿過γ?相，促進了合金的加工硬化 Zhou等[13]報道，在高鎢高溫合金的高溫拉伸過程中析出的納米γ?相顆粒有助于提高其高溫強度 Xie等[14]指出，γ?相在蠕變過程中發(fā)生定向粗化，位錯在γ-γ?界面形成密集的網(wǎng)狀結(jié)構(gòu) 這些研究結(jié)果表明，形貌和尺寸不同的γ?相在合金變形過程中與位錯的作用機制也不同 雖然對γ?相與位錯之間的關(guān)系研究得較多，但是在高溫合金的服役過程中γ?相對位錯遷移機制的影響仍不明確 研究表明，在單晶高溫合金的超溫蠕變過程中析出的二次γ?使其蠕變壽命降低[15] 而Guo等[16]發(fā)現(xiàn)，在K465合金的熱循環(huán)蠕變過程中析出的二次γ?相有助于阻礙位錯運動而使起抵抗蠕變變形的能力提高

K416B合金是典型的鎳基高鎢高溫合金，其鎢含量高達16%(質(zhì)量分數(shù))，使用溫度超過1000℃ K416B是沉淀強化型高溫合金，其中γ?相的體積分數(shù)可達60%~70%[17, 18] 因此，γ?相對K416B合金的性能至關(guān)重要 以往關(guān)于K416B合金的報道主要集中在長期時效和測試條件對其性能的影響[7, 19, 20]，對合金在高溫下析出的二次γ?相與性能的關(guān)系研究較少 K416B合金在冷卻過程中極易在γ通道中析出大量尺寸約為100 nm的細小二次γ?相，對合金的力學性能有重要的影響[21] 本文通過熱處理調(diào)整合金中γ?相的尺寸及形貌，研究其對合金拉伸及持久性能的影響

1 實驗方法

用VIM-F25型真空感應(yīng)熔煉爐熔煉實驗用K416B母合金錠，其主要化學成分(質(zhì)量分數(shù)，%)為：C 0.1、Co 7.8、Cr 5.1、W 15.8、Al 5.8、Ti 0.8、Nb 1.7、Hf 1.1、Ni其余

將母合金錠澆鑄成型，制備等軸晶試棒 從鑄態(tài)K416B合金試棒上切取直徑為3 mm、厚度為2.5 mm的圓柱形試樣，使用STA449F3型差熱分析儀進行差示掃描量熱分析(DSC)，升降溫速率為10℃/min

為了調(diào)整合金中析出相的形貌，將合金在1240℃固溶30 min后空冷，再在1160℃時效3 h后空冷，使K416B合金中保留部分大尺寸γ?相同時在γ相通道中析出細小的二次γ?相

將鑄態(tài)和熱處理態(tài)的K416B合金棒加工成標準拉伸(GB145-2001)和持久(GB/T 2039-2012)樣品(圖1) 標準拉伸樣品的工作段長度為35 mm，直徑為5 mm；標準持久樣品的工作段長度為25 mm，直徑為5 mm 用TSE504D型萬能試驗機進行拉伸測試 用F-25型蠕變/持久試驗機進行持久測試，測試條件為975℃/235 MPa，用引伸計記錄合金的變形

圖1



圖1拉伸試樣和持久試樣的加工示意圖

Fig.1Schematic diagram of mechanical test samples (a) tensile (b) durable (unit: mm)

將鑄態(tài)、熱處理態(tài)及其斷裂后的縱剖面樣品機械研磨、拋光后，用腐蝕劑5 g CuSO4+25 mL HCl+20 mL H2O進行化學腐蝕 用Stemi 508型光學顯微鏡(OM)和ZEISS MERLIN Compact場發(fā)射掃描電鏡(SEM)觀察不同狀態(tài)合金樣品的組織形貌 在距離斷口5 mm處橫向切取500 μm厚的透射電鏡樣品薄片，將其雙面機械研磨至50 μm以下后進行沖孔得到直徑為3 mm的圓片 將透射樣品在-25℃進行雙噴電解減薄，雙噴液是體積分數(shù)為10%高氯酸乙醇溶液 用Tecnai G220透射電鏡進行組織觀察和位錯分析

2 實驗結(jié)果2.1 鑄態(tài)合金的組織

圖2給出了鑄態(tài)K416B合金的晶粒和共晶組織的形貌 可以看出，合金的晶粒尺寸較大，平均尺寸約為3.5 mm 如圖2b中的箭頭所示，合金的枝晶間有大量的γ/γ?共晶，是其在凝固過程中元素偏析所致 γ/γ?共晶是高溫合金中的非平衡組織，對其高溫力學性能有不利的影響 圖3給出了K416B合金的升溫和降溫過程DSC曲線 升溫曲線上的四個特征溫度分別對應(yīng)合金中γ?相的溶解溫度，γ/γ?共晶熔化溫度，碳化物回溶溫度和合金完全熔化溫度(圖3a) 在冷卻曲線上出現(xiàn)了五個特征溫度(圖3b)，在1147℃出現(xiàn)一個微弱的放熱峰，是二次γ?相的生成溫度[21] 二次γ?相，是在高溫合金連續(xù)冷卻過程中γ?相在基體中多批次形核析出的 二次γ?相屬于后析出的γ?相，析出溫度較低，在γ通道中形核但是難以長大，因此其尺寸比初生γ?相的小[22] 按照圖4給出的合金固溶和時效熱處理理制度，得到不同于鑄態(tài)組織的、具有兩種γ?尺寸的K416B合金組織(圖5) 從圖5a可見，鑄態(tài)K416B合金中γ?相在枝晶干分布均勻，呈立方狀，平均尺寸約為200 nm 圖5b給出了K416B合金熱處理后組織形貌，可見γ?相尺寸迅速增大到1 μm左右，且其形狀由立方狀轉(zhuǎn)變?yōu)殚L條狀 另外，γ基體通道明顯粗化，細小的二次γ?相在γ基體通道中均勻析出

圖2



圖2鑄態(tài)K416B合金的OM圖像

Fig.2OM micrographs of as-cast K416B alloy (a) low magnification, (b) high magnification

圖3



圖3鑄態(tài)K416B合金的DSC曲線

Fig.3DSC results of as-cast K416B (a) heating, (b) cooling

圖4



圖4熱處理過程

Fig.4Heat treatment process

圖5



圖5鑄態(tài)和熱處理態(tài)K416B合金中γ?相的組織形貌

Fig.5Morphologies of γ? phase in K416B (a) as-cast state, (b) heat treatment

2.2 K416B合金的拉伸性能

圖6給出了鑄態(tài)和熱處理態(tài)K416B合金的應(yīng)力-應(yīng)變曲線，表1列出了兩組合金的典型性能參數(shù) 可以看出，與鑄態(tài)合金相比，熱處理態(tài)K416B合金的屈服強度明顯地由936 MPa下降到775 MPa，但是其拉伸強度沒有明顯的變化，延伸率和斷面收縮率均有所增大 相比之下，熱處理態(tài)K416B合金的塑性優(yōu)于鑄態(tài)

Table 1

表1

表1K416B合金拉伸性能

Table 1Tensile properties of K416B alloy

	Rp0.2 /MPa	Rm/MPa	A/%	Z/%
As-cast	936	1054	4.0	10.0
Heat-treatment	775	1040	6.5	11.0

圖6



圖6γ?結(jié)構(gòu)不同的K416B合金的應(yīng)力-應(yīng)變曲線

Fig.6Stress-strain curves of K416B alloy with different microstructures of γ?

2.3 拉伸斷裂后的組織形貌

圖7給出了鑄態(tài)和熱處理態(tài)K416B合金拉伸斷口的SEM形貌 由圖7a可以看出，鑄態(tài)K416B合金的斷口宏觀上較為平整，在斷口的中央可見明顯的河流狀花樣 由此可以推斷，鑄態(tài)K416B合金的室溫拉伸斷裂機制是以解理斷裂為主的脆性斷裂 在高倍數(shù)SEM照片(圖7b)中清晰可見臺階狀的解理小平面，還能觀察到沿枝晶斷裂的跡象(圖7b) 而熱處理態(tài)K416B合金的斷口整體凹凸不平，呈現(xiàn)出海綿狀特征(圖7c) 圖7d表明，在斷口中不僅出現(xiàn)解理小平面，還出現(xiàn)了少量微孔和由微孔聚集成的擴展裂紋 這表明，熱處理態(tài)K416B合金的斷裂方式主要是解理和微孔聚集，表現(xiàn)為混合斷裂

圖7



圖7γ?結(jié)構(gòu)不同的K416B合金的拉伸斷口形貌

Fig.7Fracture morphologies of K416B alloy with different microstructures of γ? (a, b) as-cast state; (c, d) heat treatment

圖8給出了K416B合金斷裂后的縱截面形貌，可見合金拉伸斷裂的二次裂紋均位于枝晶間區(qū)域 鑄態(tài)K416B合金的裂紋主要沿著枝晶間γ/γ?共晶邊緣擴展，γ?強化相的阻礙使裂紋不能進入枝晶干區(qū)域(圖8a) 而圖8b顯示，熱處理態(tài)K416B合金的二次裂紋存在于大塊γ?，與鑄態(tài)相比熱處理態(tài)的二次裂紋尺寸較大，表明大塊γ?相難以阻礙裂紋的擴展

圖8



圖8γ?相形貌不同的K416B合金拉伸斷裂后的SEM照片

Fig.8SEM images of K416B alloy with different microstructures of γ? after fracture (a) as-cast state, (b) heat treatment

圖9給出了鑄態(tài)K416B合金拉伸斷裂后的TEM明場像 可以看出，室溫拉伸斷裂后位錯在外力作用下形成了位錯環(huán)(圖9a) 另外，部分位錯以位錯對的形式切割γ?相 圖9b表明，部分切入γ?的<110>超位錯分解為層錯 圖10給出了熱處理態(tài)K416B合金γ?相的STEM-EDS圖像，可見初生γ?尺寸約為500 nm，細小的二次γ?相在γ基體中析出 圖10表明，Cr和Co元素偏析于γ基體，Al元素富集于γ?相，而W元素在γ基體和γ?相中等量分布 圖11給出了熱處理態(tài)K416B合金拉伸斷裂后的TEM圖像 由圖11a可見，在室溫拉伸期間合金基體中出現(xiàn)了位錯塞積 大量全位錯切入初生γ?相，降低了合金的屈服強度 從TEM暗場相可觀察到彎曲的位錯組態(tài)(圖11b)，表明在合金變形過程中位錯以O(shè)rowan機制繞過二次γ?，并留下相應(yīng)的位錯環(huán)

圖9



圖9鑄態(tài)K416B合金拉伸斷裂后的TEM圖像

Fig.9TEM images of as-cast K416B alloy after tensile fracture

圖10



圖10熱處理態(tài)K416B合金的STEM-EDS圖像

Fig.10STEM-EDS images of heat-treatment K416B alloy

圖11



圖11熱處理態(tài)K416B合金拉伸斷裂后的TEM圖像

Fig.11TEM images of heat-treatment K416B alloy after tensile fracture

2.4 K416B合金的持久性能

圖12給出了在975℃/235 MPa條件下鑄態(tài)和熱處理態(tài)K416B合金的持久曲線 與鑄態(tài)K416B相比，熱處理態(tài)合金的持久壽命和延伸率均有所降低 從圖12可見，合金熱處理后其持久變形第二階段的速率明顯提高、時間縮短，持久壽命取決于穩(wěn)定變形階段 鑄態(tài)K416B合金的持久壽命為89 h，斷后伸長率約為3.40% 穩(wěn)態(tài)階段應(yīng)變速率約為3.48×10-4 h-1 而熱處理態(tài)合金的持久壽命僅為28 h，斷后伸長率約為2.40%，穩(wěn)態(tài)階段應(yīng)變速率提高到7.63×10-4 h-1

圖12



圖12不同組織的K416B合金在975℃/235 MPa條件下的持久曲線

Fig.12Durable rupture curves of K416B alloys with different morphologies under the condition of 975℃/235 MPa

2.5 持久斷裂組織的形貌

圖13給出了在975℃/235 MPa條件下鑄態(tài)和熱處理態(tài)K416B合金持久斷裂后的顯微組織 圖13a、b給出了鑄態(tài)K416B合金的斷口縱剖面組織形貌，可見持久裂紋在枝晶間萌生并垂直于應(yīng)力方向 圖13b給出了合金斷口處γ?相的形貌 可以看出，γ?相在持久過程中完全筏化并出現(xiàn)了拓撲倒置現(xiàn)象，如圖中箭頭所示 鑄態(tài)K416B合金的筏形結(jié)構(gòu)連續(xù)性差，分布不均勻 圖13c、d給出了熱處理態(tài)K416B合金斷口的縱剖面組織形貌 與鑄態(tài)相比，熱處理態(tài)合金的枝晶間裂紋更加寬大(圖13c) 這表明，合金熱處理后裂紋易于在枝晶間擴展 圖13d給出了合金斷口處γ?相的形貌，可見γ?筏形的完整度高于鑄態(tài)，在γ-γ?界面還析出了碳化物顆粒

圖13



圖13鑄態(tài)和熱處理態(tài)合金在975℃/235 MPa條件下斷口縱剖面的組織

Fig.13Cross-sectional view of the alloys under the condition of 975℃/235 MPa (a, b) as-cast state; (c, d) heat treatment

圖14給出了K416B合金持久斷裂后的TEM圖像 在持久變形過程中，位錯在γ基體中沿不同方向滑移并塞積在碳化物附近(圖14a) 隨著持久變形的進行基體中的位錯密度逐漸提高，沿不同方向滑移的位錯在γ-γ?界面相互作用形成網(wǎng)狀位錯纏結(jié)(圖14b) 位錯堆積在γ基體和γ?相界面產(chǎn)生應(yīng)力集中，使部分位錯切入γ?相并在γ?相中纏結(jié)，如圖14b中的箭頭所示

圖14



圖14K416B合金在975℃/235 MPa條件下斷裂后的位錯組態(tài)

Fig.14Dislocation structures of K416B after rupture under 975℃/235 MPa

圖15給出了K416B合金持久斷裂后的STEM-EDS圖像 從圖15a可見，在K416B合金的持久變形過程中大量納米級碳化物顆粒在基體通道內(nèi)析出，在圖的右下角可見少量碳化物出現(xiàn)在γ?相中 圖15b給出了EDS面掃結(jié)果，可見在碳化物中富集W元素，由此判定碳化物為M6C型 還可以看出，大量碳化物在兩相界面析出，延緩了γ?相的筏化速度，使局部γ?相仍能維持一定的立方度

圖15



圖15K416B合金在975℃/235 MPa條件下蠕變斷裂后的STEM明場相和對應(yīng)的能譜

Fig.15STEM bright field image (a) and corresponding elemental mapping (b) of K416B after creep rupture under 975℃/235 MPa

3 討論3.1 γ? 相對K416B合金拉伸性能的影響

K416B合金中γ?相，對合金的拉伸斷裂方式有重要的影響 在合金的室溫拉伸過程中，應(yīng)力傾向于集中在枝晶間區(qū)域 枝晶間的碳化物其強度較高，在拉伸過程中破碎產(chǎn)生裂紋[23] 隨后，裂紋在相鄰的共晶中擴展 與大塊γ?相不同，γ/γ?共晶中的γ通道能改變裂紋擴展方向，使裂紋沿共晶邊緣生長[24] 因此，鑄態(tài)K416B合金發(fā)生脆性斷裂 而熱處理后共晶中的γ相溶解，大塊γ?相發(fā)生更大的塑性變形，裂紋在枝晶間區(qū)域的擴展加快 隨著拉應(yīng)力的增大，裂紋在大塊γ?相中形成的微孔洞聚集連接并沿晶界擴展 合金熱處理后枝晶間的γ/γ?共晶轉(zhuǎn)變?yōu)棣?相，是合金發(fā)生混合斷裂的主要原因

在含有高體積分數(shù)沉淀強化相的高溫合金中，位錯切過γ?的臨界應(yīng)力小于Orowan繞過應(yīng)力 研究表明，在奧氏體高溫合金中位錯通過γ?的方式與γ?尺寸相關(guān)[25] 在臨界尺寸以下，位錯以位錯對的形式切入γ?相 屈服強度位于刃型位錯對和螺型位錯對切割γ?的理論強度之間，并隨著γ?相尺寸的增大而提高 γ?相的尺寸超過臨界尺寸后位錯以O(shè)rowan機制繞過γ?強化相，隨著γ?相之間距離的增大屈服強度的提高緩慢 本文的實驗結(jié)果表明，在鑄態(tài)K416B合金拉伸期間形變位錯在基體通道中滑移，γ?強化相有效阻礙了位錯運動 位錯塞積在γ-γ?界面導致應(yīng)力集中，少量位錯以位錯對的形式切入γ?相 此外，K416B合金中W元素的含量較高使合金中γ?相的層錯能降低，<110>超位錯在γ?相中分解形成層錯(圖9b) 熱處理后K416B合金的初生γ?迅速粗化 與鑄態(tài)K416B合金相比，在熱處理態(tài)合金的拉伸變形過程中更多位錯切入γ?，使合金產(chǎn)生更大的塑性變形 這表明，γ?相的強度隨著其尺寸的增大而降低，阻礙位錯運動的能力減弱 因此，熱處理后K416B合金強度降低而延伸率提高

值得注意的是，與先前研究中具有大尺寸γ?相的K416B合金(710 MPa/5.0%)相比[19, 20]，本文通過熱處理制備的K416B合金(775 MPa/6.5%)其屈服強度和延伸率更高 其原因是，在熱處理態(tài)K416B合金中除了初生γ?相，在基體中還析出了大量二次γ?相 二次γ?相的生成降低了γ通道的寬度，增大了位錯切割γ?的臨界切應(yīng)力，使變形過程中位錯以O(shè)rowan機制繞過小尺寸二次γ?相 這表明，二次γ?反向疇界能較高、硬度較大，位錯難以通過切割的方式進入γ?相

3.2 γ? 相對K416B合金持久變形的影響

在持久變形過程中，合金中的位錯密度逐漸提高 由于γ?相的強度高于γ基體，基體的塑性變形量高于γ?相，使γ基體和γ?強化相界面處產(chǎn)生應(yīng)力集中并出現(xiàn)密集的位錯網(wǎng)絡(luò)(圖14) 在持久變形期間，合金中的γ?相發(fā)生粗化 鑄態(tài)K416B合金中的γ?相在基體中均勻分布，發(fā)生筏化所需的時間更長，位錯通過γ?相需要更多的能量，因此鑄態(tài)K416B合金的持久壽命可達90 h 與之相比，熱處理態(tài)K416B合金的持久壽命僅為30 h，其大幅度下降可歸因于：1)熱處理后γ?相的尺寸增大，γ基體通道的寬度也隨之增大 在γ通道中析出了細小的二次γ?相和元素擴散距離減小，使γ?的筏化速率提高和合金抵抗變形的能力減弱 2)熱處理后初生γ?相尺寸的增大促進位錯切入初生γ?相，并提高了持久應(yīng)變速率(圖11) 因此可以觀察到，熱處理態(tài)K416B合金持久斷裂后γ?筏形更加完整(圖13d) 形變位錯切入初生γ?相誘導顆粒狀W6C在γ?相中彌散析出(圖15)消耗了大量的W元素，使γ-γ?兩相錯配度減小，降低了合金的強化水平，導致持久壽命大幅度降低[26~28]

4 結(jié)論

(1) K416B合金中γ?相尺寸的增大使其屈服強度降低和延伸率提高 小尺寸二次γ?相有助于提高合金的屈服強度，大尺寸初生γ?相使其塑性提高 K416B合金的持久壽命隨γ?相尺寸的增大而下降

(2) 在K416B合金的拉伸過程中產(chǎn)生的裂紋主要萌生于枝晶間區(qū)域 在鑄態(tài)K416B合金的拉伸變形期間基體中形成高密度位錯纏結(jié)，單根位錯切入γ?相并分解為層錯，另有部分位錯以位錯對的形式切入γ?相 在熱處理態(tài)K416B合金中，在基體通道內(nèi)滑移的位錯受到二次γ?的阻礙，位錯以位錯對或全位錯的形式切入大尺寸γ?相，以O(shè)rowan機制繞過二次γ?相

(3) 在熱處理后的K416B合金中初生γ?相的尺寸增大，使持久變形期間更多的位錯切入γ?相并促進顆粒狀W6C析出 在γ通道中析出的二次γ?相提高了合金的筏化速率，使熱處理態(tài)K416B合金的持久壽命低于鑄態(tài)

參考文獻

View Option 原文順序文獻年度倒序文中引用次數(shù)倒序被引期刊影響因子

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通過OM、SEM和XRD對高W鎳基高溫合金進行組織觀察與分析，研究了W含量對鎳基高溫合金凝固組織的影響規(guī)律 結(jié)果表明，當W含量為14% (質(zhì)量分數(shù)，下同)時，鎳基合金中無α-W相析出 當W含量高于16%時，合金凝固期間可析出α-W，并且隨W含量提高，合金的晶粒尺寸由1.04 mm減小至0.17 mm，共晶含量由6%增至10%；W含量對在枝晶間/枝晶干內(nèi)的γ'相尺寸及形態(tài)無明顯影響 由于α-W的析出溫度較高，在凝固期間首先析出，并在殘余液相收縮作用下，α-W向液相核心處發(fā)生轉(zhuǎn)移并長大；同時α-W可作為異質(zhì)形核的核心，降低枝晶形核的臨界形核功，使18%W合金獲得較小的晶粒尺寸 此外，在不同取向枝晶匯聚生長的作用下，殘余液相中Al、Ti等元素形成較高的濃度梯度而發(fā)生共晶轉(zhuǎn)變，這是18%W合金中共晶含量較高的主要原因

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Differential scanning calorimetry (DSC) analysis, isothermal solidification experiment and Thermo-Calc simulation were employed to investigate solidification characteristics of K417G Ni-base superalloy. Electron probe microanalysis (EPMA) was employed to analyze the segregation characteristics. Liquidus, solidus and the formation temperatures of main phases were measured. In the process of solidification, the volume fraction of liquid dropped dramatically in the initial stage, while the dropping rate became very low in the final stage due to severe segregation of positive segregation elements into the residual liquid. The solidification began with the formation of primary γ. Then with solidification proceeding, Ti and Mo were enriched in the liquid interdendrite, which resulted in the precipitation of MC carbides in the interdendrite. Al accumulated into liquid at the initial stage, but gathered to solid later due to the precipitation of γ/γ′ eutectic at the intermediate stage of solidification. However, Co tended to segregate toward the solid phase. In the case of K417G alloy, combining DSC analysis and isothermal solidification experiment is a good way to investigate the solidification characteristics. Thermo-Calc simulation can serve as reference to investigate K417G alloy.

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通過OM、SEM和XRD對高W鎳基高溫合金進行組織觀察與分析，研究了W含量對鎳基高溫合金凝固組織的影響規(guī)律 結(jié)果表明，當W含量為14% (質(zhì)量分數(shù)，下同)時，鎳基合金中無α-W相析出 當W含量高于16%時，合金凝固期間可析出α-W，并且隨W含量提高，合金的晶粒尺寸由1.04 mm減小至0.17 mm，共晶含量由6%增至10%；W含量對在枝晶間/枝晶干內(nèi)的γ'相尺寸及形態(tài)無明顯影響 由于α-W的析出溫度較高，在凝固期間首先析出，并在殘余液相收縮作用下，α-W向液相核心處發(fā)生轉(zhuǎn)移并長大；同時α-W可作為異質(zhì)形核的核心，降低枝晶形核的臨界形核功，使18%W合金獲得較小的晶粒尺寸 此外，在不同取向枝晶匯聚生長的作用下，殘余液相中Al、Ti等元素形成較高的濃度梯度而發(fā)生共晶轉(zhuǎn)變，這是18%W合金中共晶含量較高的主要原因

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[J]. Acta Metall. Sin., 2015, 51(8): 943

[本文引用: 2]

謝 君, 于金江, 孫曉峰等 溫度對高W含量K 416

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Chen M K, Xie J, Shu D L, et al.

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Tensile, creep behavior and microstructure evolution of an as-cast Ni-based K417G polycrystalline superalloy

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The Ni-based K417G superalloy is extensively applied as aeroengine components for its low cost and good mid-temperature (600-900 °C) properties. Since used in as-cast state, the comprehensive understanding on its mechanical properties and microstructure evolution is necessary. In the present research, the tensile, creep behavior and microstructure evolution of the as-cast K417G superalloy under different conditions were investigated. The results exhibit that tensile cracks tend to initiate at MC carbide and γ/γ′ eutectic structure and then propagate along grain boundary. As the temperature for tensile tests increases from 21 °C to 700 °C, the yield strength and ultimate tensile strength of K417G superalloy decreases slightly, while the elongation to failure decreases greatly because of the intermediate temperature embrittlement. When the temperature rises to 900 °C, the yield strength and ultimate tensile strength would decrease significantly. The creep deformation mechanism varies under different testing conditions. At 760 °C/645 MPa, the creep cracks initiate at MC carbides and γ/γ′ eutectic structures, and propagate transgranularly. While at 900 °C/315 MPa and 950 °C/235 MPa, the creep cracks initiate at grain boundary and propagate intergranularly. As the creep condition changes from 760 °C/645 MPa to 900 °C/315 MPa and 950 °C/235 MPa, the γ′ phase starts to raft, which reduces the creep deformation resistance and increases the steady-state deformation rate.

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