7050鋁合金採用的階段式淬火及時效(Step-Quench and Aging, SQA)熱處理，無論是熱輥軋材、冷輥軋材以及中間熱機處理(ITMT)材，其抗SCC的能力均有顯著地改善效果，不但大幅超越了RRA與T6處理，也較T73處理為佳，同時在強度上也大於T73處理者。對7050鋁合金而言，較高的再結晶比例會造成不整合(incoherent)分散粒子Al3Zr及高角度晶界數目的增加，因而提高材料的淬火敏感性，而導致強度下降。此外，高角度晶界比例的增加，也會提高SCC敏感性。冷輥軋材有較高的再結晶比例，因此強度及抗SCC能力較低。相比較下，熱輥軋材可獲得最少的再結晶比例，因而有最高的強度及優異的抗SCC能力。ITMT製程能產生細晶(grain refinement)的組織，可提高滑移的均勻性，因而提高抗SCC的能力，然而在強度上未能有效提昇。對7050合金而言，在高溫固溶處理後實施階段式淬火及時效(SQA)熱處理，可使晶界析出物粗大化及間距加大，且晶內析出緻密的強化相，因而可有效提昇抗SCC的能力並能兼顧到強度的要求。在本研究中若採用熱輥軋製程並實施SQA(220℃?10s or 200℃?30s)熱處理法，不但可得到比熱輥軋之T73處理還高的強度及抗SCC能力，同時也大幅減少時效處理的時間。
對7050及7075合金而言，再結晶組織及淬火敏感性，明顯的受到均質化條件所造成的分散粒子分佈的影響。當均質化處理的溫度愈高，其分散粒子分佈愈粗大稀疏，熱輥軋、固溶處理後再結晶的比例會愈高；而施以階段式均質化條件(Step-Homogenization, Step-H)，可獲得最微細且緻密的分散粒子分佈，在熱輥軋、固溶處理時抑制再結晶的能力最大，再結晶的比例為最低。對含Zr之7050合金而言，細密的分散粒子分佈，可降低其淬火敏感性，但對含Cr之7075合金而言，反而會提高其淬火敏感性。對7050合金而言，施以均質化條件Step-H處理的試片經時效處理後其強度均較其他均質化條件處理之試片為高；而施以階段淬火及時效熱處理法(SQA)之試片在再結晶比例小於70%時其強度都高於T73處理者。但對於7075合金，施以SQA熱處理法之強度均遠低於T73處理者。7050合金在固溶處理時採用階段淬火及時效熱處理法(SQA)，可有效提升其抗應力腐蝕破壞性並且兼顧到強度。且再結晶比例越低，其強度及抗應力腐蝕破壞性也越高。在本研究中，7050合金經Step-H階段式均質化處理並配合在固溶處理後採用SQA時效熱處理法，可獲得最高抗應力腐蝕破壞性與較佳的抗拉強度，且均較工業應用之傳統T73製程優異，實為一能有效兼顧抗應力腐蝕破壞性與強度之最佳製程。然而此製程對7075合金而言，則因淬火敏感性高而不適用。
7475-T7351鋁合金銲接熱影響區(HAZ)強度降低的主要原因為η?相轉變成穩定相η及η相之粗化所造成。而銲後施以T73人工時效(PWAA-T73)並不能提昇熱影響區強度。熱影響區之衝擊韌性與降伏強度有相反之趨勢，η相愈多且愈粗化韌性就愈高。7475-T7351縱向試片(7475L)的熱影響區強度、延性及韌性，均比短橫軸向試片(7475ST)來得高。
7005-T1銲接熱影響區(HAZ)強度弱化的原因是由於銲接熱循環造成差排的消解(annihilation)及再結晶所致。此HAZ之強度可藉由銲後T53人工時效(PWAA-T53)來提昇。相對而言，7005-T53銲接熱影響區強度弱化的主要原因是由於η?相的粗化及轉變成穩定相η所造成。而強度最低值出現在峰值溫度為250℃之熱循環附近，此過時效區域的強度無法藉由銲後T53人工時效來恢復。而在峰值溫度高於326℃以上之熱影響區域，因強化相η?之回溶造成強度弱化，此區域強度則可藉由銲後T53人工時效恢復至銲接前之強度。若考量HAZ之強度並於銲後施以T53人工時效，則選擇以T1狀態銲接可獲得較T53銲後為佳的強度。對T1狀態及T53狀態之7005合金而言，銲後施以T53人工時效均可提昇其熱影響區之衝擊韌性，而η?的粗化或析出穩定相η可提高其衝擊韌性，但會使強度(UTS及YS)降低。
7475鋁合金比7005鋁合金有較低的高溫強度、較差的熱延性與延性恢復能力，以及較寬的脆性溫度範圍(BTR)，因而有較高的熱裂敏感性。而7475鋁合金之短橫軸向試片(7475ST)較縱向試片(7475L)有較高的熱裂敏感性。對以上兩合金而言，在熱延性試驗中，從零強度溫度(NST)冷卻過程所得到的強度及熱延性均小於從零延性溫度(NDT)冷卻過程所獲得的，此表示承受較高峰值溫度之HAZ區域會有較高的熱裂敏感性。7475鋁合金熱影響區熱裂發生的主要原因是由於低熔點之晶界偏析物或共晶相在高溫下產生不平衡共晶熔解而導致晶界液化(liquation)所致。而這些偏析物或共晶相含有較高含量的Mg、Cu、Zn之溶質元素。且由掃瞄式電子顯微鏡(SEM)之破斷面觀察及分析發現7475合金隨著延性之降低，破裂型態由延性穿晶轉為脆性沿晶模式。

For the 7050 alloy, the higher fraction of recrystallization would cause a larger amount of incoherent Al3Zr dispersoids and high-angle grain boundaries, which increases quench sensitivity and SCC susceptibility of the alloy. Therefore, both strength and SCC resistance of the cold-rolled alloy are low. On the other hand, the hot-rolled alloy has the highest strength and superior SCC resistance, as a result of the lowest fraction of recrystallization. The SCC resistance can be significantly enhanced owing to the fine grain structure obtained with the intermediate thermomechanical treatment (ITMT); however, the strength cannot be effectively improved. The proposed treatment, step-quench and aging (SQA), can significantly improve SCC resistance and attain optimum strength by controlling the grain boundary and matrix microstructures. Higher SCC resistance and strength together with much reduced aging time are obtained simultaneously with the proposed SQA(220℃/10s or 200℃/30s) after hot-rolling, as compared to those with the conventional T73 treatment.
The recrystallized structure and the quench sensitivity of the 7050 and 7075 aluminum alloys were found to be significantly affected by dispersoid distribution depending on the homogenization conditions. The finest and densest dispersoid distribution, generated through a step-homogenization (Step-H) treatment, can effectively inhibit recrystallization to obtain the smallest fraction of recrystallized structure. The distribution considerably lowers the quench sensitivity of the 7050 alloy, but increases the quench sensitivity of the 7075 alloy. For the 7050 alloy, the Step-H always achieved the highest strength and SCC resistance in all aged temper. The proposed SQA treatment can effectively improve both the SCC resistance and strength of 7050 alloy. In particular, combining the SQA treatment with the Step-H can result in much greater strength and SCC resistance than can be achieved by the conventional T73 treatment. However, this SQA treatment is not applicable to the 7075 alloy because of its high quench sensitivity.
The microstructure and mechanical properties of the HAZ of the 7475 and 7005 alloys were considerably influenced by the peak temperatures of thermal cycles. The decay of strength in weld HAZ of the 7475-T7351 alloy was primarily due to the η? phase transformation to the η phase and coarsening of the η phase. The HAZ strength in overaged zone cannot be improved by post-weld T73 artificial aging (PWAA-T73) treatment. The HAZ toughness of the 7475-T7351 alloy had an opposite trend as the yield strength. The HAZ strength, toughness, and ductility of the 7475-T7351 alloy in longitudinal direction were always greater than those in short-transverse direction.
The decay of strength in weld HAZ of the 7005-T1 alloy is due to the annihilation of dislocations and recrystallization. In contrast, for the 7005-T53 alloy, The decay of HAZ strength is primarily caused by the coarsening of theη? phase and transformation to the η phase. The yield strength of the HAZ in the 7005-T1 alloy is significantly improved by the postweld T53 aging (PWAA-T53) treatment. However, the 7005-T53 alloy has the weakest HAZ region at peak temperature 250℃. This overaged zone has the maximum amount of the coarser η? and ηphases, and does not respond well to postweld T53 aging. For peak temperature at 326℃ and above, the strength and toughness curves of the HAZ in both T1 and T53 tempered 7005 alloys coincide because of the same microstructure. The postweld T53 aging treatment can improve the impact toughness of the HAZ.
As compared to the 7005 alloy, the 7475 alloy exhibited lower hot ductility and strength, significantly poor ductility recovery on cooling from the nil-strength temperature (NST), considerably wider brittle temperature range (BTR) and crack susceptible region (CSR), and thus greater susceptibility to HAZ hot cracking. The 7475 alloy in longitudinal direction had a lower crack susceptibility than that in short-transverse direction. The results showed good agreement with those of the spot-Varestraint tests. Scanning electron microscopy (SEM) revealed that the ductility loss was accompanied by fracture transition from ductile transgranular mode to brittle intergranular mode. The HAZ hot cracking in alloy 7475 due to the loss of ductility was primarily caused by liquation of low-melting-point grain boundary segregates or eutectics that contained great amount of Mg, Cu, and Zn.

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